Nitride semiconductor heterostructures and related methods

ABSTRACT

Semiconductor structures and devices based thereon include an aluminum nitride single-crystal substrate and at least one layer epitaxially grown thereover. The epitaxial layer may comprise at least one of AlN, GaN, InN, or any binary or tertiary alloy combination thereof, and have an average dislocation density within the semiconductor heterostructure is less than about 10 6  cm −2 .

RELATED APPLICATIONS

This application is a continuation of U.S. patent application Ser. No.11/503,660, filed Aug. 14, 2006, which is a continuation-in-part of U.S.patent application Ser. No. 11/431,090, filed May 9, 2006, which is acontinuation-in-part of U.S. patent application Ser. No. 10/910,162,filed Aug. 3, 2004, which is itself a continuation-in-part of U.S.patent application Ser. No. 10/324,998, filed Dec. 20, 2002, issued asU.S. Pat. No. 6,770,135 on Aug. 3, 2004, which claims the benefit of andpriority to U.S. Provisional Application Ser. No. 60/344,672, filed Dec.24, 2001. The entire disclosures of each of these applications areincorporated herein by reference.

GOVERNMENT SPONSORSHIP

The U.S. Government may have certain rights in this invention pursuantto SBIR Contract N00014-98-C-0053 awarded by the Office of NavalResearch, under SBIR Contract F33615-00-C-5531 awarded by the Air ForceResearch Laboratory, under SBIR Contract F33615-00-C-5425 awarded by theMissile Defense Agency and under ATP Contract 70NANB4H3051 awarded bythe National Institute of Standards and Technology (NIST).

FIELD OF THE INVENTION

The present invention relates to the fabrication of single-crystalaluminum nitride (AlN) structures, and, more particularly, to growingrelatively large, single-crystal AlN substrates bysublimation-recondensation technique at growth rates exceeding 0.3mm/hr.

BACKGROUND OF THE INVENTION

Several types of materials are routinely used to fabricate substratesfor nitride-based semiconductor structures, which, in turn, can beemployed as components of high-performance electronic and optoelectronicdevices.

For devices employing GaN or Ga_(1-x)In_(x)N, the most desirablesubstrate would nominally be a large-area GaN single-crystal wafer.While several methods to grow GaN crystals have been proposed, none ofthem appears to be commercially feasible to fabricate large-area bulkcrystals of GaN.

Sapphire is a popular conventional substrate material, becauserelatively high-quality, inexpensive sapphire substrates arecommercially available. However, sapphire is far from being perfectlysuited for GaN epitaxy. Its lattice mismatch to GaN is relatively large(about 16%), it has little distinction between the + and − [0001]directions which can give rise to +/−c-axis domains in epitaxial filmsof GaN, and its differential thermal expansion may lead to crackingduring the post-fabrication cooling process. In spite of thosedrawbacks, recently, Nichia Ltd of Japan has announced the production ofa violet laser with significant commercial possibilities—more than10,000 hours of operating life using sapphire substrates. Currently,laser diodes (“LDs”) are relatively expensive. Using sapphire substratesleads to a costly fabrication process because it requires growing bufferlayers and using lateral epitaxial overgrowth techniques. Even thoughNichia's announcement is very promising, violet lasers grown oversapphire substrates still have shortcomings. For example, heat may buildup in these lasers during operation. Sapphire, with its very low thermalconductivity, traps that heat, which may trigger a premature burnout ofthese devices. To build a more durable blue laser, Nichia and others areinvestigating other alternatives, such as free-standing GaN substrates.In this technique, the substrate is removed after a thick GaN layer isgrown thereover. This method leaves the GaN as the base for building thelaser. This base should be better at dissipating heat, in addition tomatching the alloy layers disposed thereover. However, this approachfurther increases process complexity and associated fabrication costs.

Single-crystal substrates of SiC present an attractive alternative tosapphire due to their close lattice match to AlN/GaN in the planeperpendicular to the c-axis (the so-called “c-plane”) and high thermalconductivity. In addition, SiC substrates can be made electricallyconducting, which is attractive for some applications, includingfabrication of LEDs and LDs. However, wurtzite SiC (matching thewurtzite crystal structure of GaN) is not available and the latticemismatch along the c-axis between GaN and both 4H and 6H SiC issubstantial. In addition, the chemical bonding between the Group IVelements of the SiC and the Group III or Group V elements of thenitrides is likely to create nucleation problems leading to electronicstates at the interface.

It is, therefore, desirable to provide alternative substrates forcommercial fabrication of nitride-based (e.g., GaN) devices. Inparticular, the physical and electronic properties of AlN—wide energybandgap (6.2 electron volts), high breakdown electric field, extremelyhigh thermal conductivity, and low optical density—afford this materiala great potential for a wide variety of semiconductor applications as asubstrate material. Despite these useful properties of AlN substrates,commercial feasibility of AlN-based semiconductor devices, however, havebeen limited by the unavailability of large, low-defect single crystalsof AlN.

A sublimation-recondensation technique was developed for AlN crystalgrowth by Slack and McNelly (G. A. Slack and T. McNelly, J. Cryst.Growth 34, 263 (1976) and 42, 560 (1977), hereinafter the “Slackreference,” incorporated herein by reference). In this technique,polycrystalline source material is placed in the hot end of a cruciblewhile the other end is kept cooler. The crystal nucleates in the tip andgrows as the crucible is moved through the temperature gradient. Thisapproach demonstrated relatively slow crystal growth of 0.3 mm/hr whilethe crystal growth chamber was maintained at 1 atm of N₂, producing aconical crystal about 12 mm in length and about 4 mm in diameter.

To make AlN substrates and devices built thereon commercially feasible,it would be desirable to increase the growth rate. A number ofresearchers have examined the possibility of such increase. Many ofthem, however, relied on rate equations derived by Dryburgh (see“Estimation of maximum growth rate for aluminum nitride crystals bydirect sublimation,” J. Crystal Growth 125, 65 (1992)), which appear tooverestimate the growth rate of AlN and, in particular, suggest that themaximum growth conditions are near stoichiometric vapor conditions,i.e., the Al and N₂ partial pressures should be adjusted so that the Alpartial pressure is twice that of the N₂. See, for example, U.S. Pat.Nos. 5,858,085; 5,972,109; 6,045,612; 6,048,813; 6,063,185; 6,086,672;and 6,296,956, all to Hunter. In addition, known approaches generallymaintain the N₂ partial pressure at less than atmospheric pressure.

Most attempts at increasing the growth rate of AlN crystals under suchstoichiometric and/or sub-atmospheric pressure conditions have met withlimited success. In addition, it appears to be impossible to achieve thegrowth rate, or the electronics-grade quality Hunter discloses in hispatents with the N₂ pressure below one atmosphere. Additional researchin the field of fabrication of AlN crystals was reported by Segal andhis colleagues. (See Segal et al. “On mechanisms of sublimation growthof AlN bulk crystals,” J. Crystal Growth 211, 68 (2000)). This articleappears to be the first peer-reviewed publication suggesting thatDryburgh's growth equations are incorrect. Segal and his colleagues,however, teach open growth conditions, allowing the Al vapor to escape.Disadvantageously, it would be difficult to grow large boules of AlNusing this approach, because: (i) due to the non-uniform growth acrossthe surface, growth control is difficult; (ii) a large amount of Alwould be wasted; (iii) the excess Al in the rest of the furnace wouldcreate problems due to of its high reactivity; and (iv) it would bedifficult to maintain a high temperature differential between the sourceand growing crystal surface.

Critical to the ability to made commercially practical nitride-basedsemiconductor devices, is achieving low levels of defect densitiesthroughout the semiconductor structure. With conventional nitridesemiconductor growth techniques on commonly available foreign substratesor high-defect nitride substrates, a GaN buffer layer is grown thick toachieve planar growth fronts and to relax the GaN epitaxial layers priorto formation of the active region of the device, i.e. the GaN/AlGaNheterostructure. This approach will result in epitaxial layers withdefect densities in the 10⁸ cm⁻² to 10¹⁰ cm⁻² range.

Other researchers have attempted to reduce dislocation densities byepitaxial lateral overgrowth, novel nucleation schemes to initiategrowth on the foreign substrates, or by adding complex structure such assuperlattices into the epitaxial profile. For example, Sumitomo Electrichas reported local regions of 10⁵ dislocations per cm² although theaverage dislocation density in their GaN substrates exceeds 10⁶ cm⁻².Also, researchers from Tokyo University of Agriculture and Technologyand TDI reported the defect densities of 10⁷ cm⁻². Work by others on AlNtemplates can be expected to have a large defect density due to theinitial growth on a foreign substrate, even though the originalsubstrate might be removed to obtain a freestanding AlN wafer.

Thus, present state-of-the-art technology employs foreign substratesthat have large thermal and lattice differences relative to AlN, AlGaN,and GaN epitaxial layers. This results in defect densities ranging from10⁸ cm⁻² to 10¹⁰ cm⁻², which makes certain devices—especially devicesemploying AlGaN layers with high Al content—unrealizable.

A need therefore exists for large AlN substrates suitable forfabricating semiconductor devices thereon and commercially-feasiblemethods for manufacturing these substrates that address theaforementioned drawbacks.

SUMMARY OF THE INVENTION

Accordingly, in its various aspects and embodiments, the presentinvention focuses on methods and apparatus for producing AlN substratesthat advantageously have a relatively small lattice mismatch (around2.2%) with respect to GaN, and which have an almost identical thermalexpansion from room temperature to 1000° C.

Furthermore, bulk AlN crystals grown using various embodiments of theclaimed fabrication methods may have a radial dimension exceeding 10-15mm and a length exceeding 5 mm. Also, the claimed invention facilitatesfabrication of AlN crystals having very low dislocation densities—under10,000 cm⁻², typically, about 1,000 cm⁻² or less, and, in someembodiments, being substantially devoid of dislocation defects. Thesebulk crystals, in turn, enable the fabrication of high-quality AlNsubstrates having surfaces of any desired crystallographic orientationby slicing them out of properly oriented bulk crystals. Possibleorientations include the c-face which is cut parallel to the (0001)plane, the a-face which is cut parallel to the (1120) plane, and them-face which is cut parallel to the (1010) plane.

These AlN crystals also advantageously have the same wurtzite crystalstructure as GaN and nominally the same piezoelectric polarity. Also,the chemical compatibility with GaN is much better than that of SiC. Inaddition, AlN substrates cut from the bulk crystals tend to beattractive for Al_(x)Ga_(1-x)N devices requiring higher Al concentration(e.g., for high-temperature, high-power, radiation-hardened, andultra-violet wavelength applications). Emerging optoelectronic andelectronic devices based on the epitaxial growth of GaN andAl_(x)Ga_(1-x)N layers may significantly benefit from such an improvedsubstrate.

Generally, in one aspect, the invention features an apparatus forgrowing a bulk single-crystal aluminum nitride. The apparatus includes ahousing defining a growth chamber, the housing including a gas outletconfigured for selectively evacuating and venting the growth chamber, agas inlet configured for pressurizing the growth chamber, and a viewingport configured for pyrometric monitoring of crystal growth temperatureswithin the growth chamber. A radio frequency (“RF”) coil is disposedwithin the growth chamber and configured for inducing an electromagneticfield therein. A quartz tube is disposed coaxially within the coil. Afirst set of shielding is disposed coaxially within the quartz tube,including from about 5 to about 7 concentric pyrolytic boron nitride(pBN) cylinders, each of the pBN cylinders having a wall thickness ofgreater than about 0.05 inches (0.13 cm), each of the cylinders having alength dimension along the longitudinal axis greater than the lengthdimension of the coil. A second set of shielding is disposed coaxiallywithin the first set of shielding, the second set of shielding includingtwo concentric, open joint tungsten cylinders, each of the tungstencylinders having a wall thickness of less than about 0.005 inches (0.013cm); each of the tungsten cylinders having a length dimension along thelongitudinal axis less than the length dimension of the RF coil. A pushtube is disposed coaxially within the second set of shielding; the pushtube having a proximal side and a distal side, the distal side includinga set of metallic baffles having a center hole which provides for thepyrometric monitoring of crystal growth temperatures, the proximal sideincluding another set of metallic baffles. A crucible is disposedcoaxially within the push tube, the crucible having a conically shapeddistal end and a proximal end; the crucible defining a crystal growthenclosure; the proximal end including a high purity, polycrystallinealuminum nitride source material, the distal end being configured forgrowth of the bulk single crystal aluminum nitride. The push tube isdisposed on a push rod assembly configured for sliding the crucible andthe push tube along the longitudinal axis. The first set of shieldingand the second set of shielding are configured to provide a thermalgradient axially within the cavity of the crucible of greater than about100° C./cm.

In another aspect, the invention is directed to a method for growingbulk single crystals of aluminum nitride. The method includes utilizingthe apparatus described above, purging the growth chamber by evacuatingthe growth chamber to a pressure less than or equal to about 0.01 mbar(1 Pa), and refilling the growth chamber with substantially purenitrogen gas to a pressure of about 1 bar (100 kPa). The growth chamberis then evacuated to a pressure less than or equal to about 0.01 mbar (IPa), and then pressurized to about 1 bar (100 kPa) with a gas includingabout 95% nitrogen and about 5% hydrogen. The chamber is heated to afirst temperature, the heating including ramping the temperature of theconical upper end of the crucible to about 1800° C. in a period of about15 minutes. The growth chamber is then pressurized to about 1.3 bar (130kPa) with the gas including about 95% nitrogen and about 5% hydrogen,and heated to a growth temperature. A distal end of the crucible is thenramped to about 2200° C. in a period of about 5 hours. The push tube andthe crucible are moved axially through the growth chamber at a rate ofabout 0.4 to about 0.9 millimeters per hour, wherein single crystals ofaluminum nitride are grown.

In yet another aspect, the invention focuses on a method for growingbulk single crystals of aluminum nitride. The method includes evacuatinga growth chamber, pressurizing the growth chamber to about 1 bar with agas including about 95% nitrogen and about 5% hydrogen, and placingsource polycrystalline AlN in a proximal end of a crystal growthenclosure. The method further includes placing a distal end of thecrystal growth enclosure in a high temperature region of the growthchamber, ramping the high temperature region to about 1800° C.,maintaining pressure in the growth chamber at about 1.3 bar, and rampingthe high temperature region to about 2200° C. The distal end of thecrystal growth enclosure is moved towards a low temperature region ofgrowth chamber at a rate of about 0.4 to about 0.9 millimeters per hour,wherein a single crystal of aluminum nitride grows at the distal end ofthe crystal growth enclosure.

A further aspect of the invention involves a method of producing bulksingle crystals of AlN, which includes providing in a crystal growthenclosure Al and N₂ vapor capable of forming bulk crystals, maintainingin the crystal growth enclosure a N₂ partial pressure which is greaterthan the stoichiometric pressure relative to the Al, maintaining thetotal vapor pressure in the crystal growth enclosure atsuper-atmospheric pressure, and providing at least one nucleation sitein the crystal growth enclosure. The method also includes cooling thenucleation site relative to other locations in the crystal growthenclosure, and depositing the vapor under conditions capable of growingsingle crystalline AlN originating at the nucleation site. A variationof this aspect includes producing a prepared substrate by cutting awafer or a cylinder from the bulk single crystal; preparing a surface onthe wafer or cylinder receptive to an epitaxial layer; and depositing anepitaxial layer or a complete crystal boule on the surface.

In yet another aspect, the invention features a system for producingbulk single crystals of AlN. The system includes a source of Al and N₂vapor, a crystal growth enclosure for containing the vapor, and at leastone nucleation site in the crystal growth enclosure. The crystal growthenclosure has a selective barrier configured to permit migration of N₂therethrough, and to substantially prevent migration of Al therethrough.A pressurization system is configured to maintain, in the crystal growthenclosure, a N₂ partial pressure greater than stoichiometric pressurerelative to the Al, and to maintain the total vapor pressure in thecrystal growth enclosure at super-atmospheric pressure. A selectiveheating system is configured to maintain the nucleation site at atemperature lower than at other locations in the crystal growthenclosure.

In yet another aspect, the invention features a single-crystal AlN boulehaving a diameter greater than about 25 mm and a dislocation densityless than about 10,000 cm⁻². In various embodiments, the thickness ofthe boule is greater than about 15 mm. In some embodiments, the AlNboule has a dislocation density of about 10,000 cm⁻² or less, forexample, 5,000 cm⁻² or less, or, more particularly, about 1,000 cm⁻² orless. In certain embodiments, the dislocation density is about 500 cm⁻²,or as low as 100 cm⁻². In a particular embodiment, the boule has atleast one region substantially devoid of dislocation defects.

Also, in still another aspect, the invention involves a method forfabricating a semiconductor device, such as a solid-state laser diode ora solid state LED. The method includes forming a bulk single crystal ofaluminum nitride using any of the methods described above, followed byremoving the bulk single crystal of aluminum nitride from the crystalgrowth enclosure and cutting it to form a substrate of AlN. Thesubstrate may have a diameter greater than about 25 mm, a thickness ofless than about 1 mm—for example, about 500 μm or about 350 μm—and adislocation density less than about 10,000 cm⁻². The method furtherincludes depositing, on the AlN substrate, at least two layersindependently selected from the group consisting of AlN, GaN, InN, andbinary or tertiary alloy combinations thereof.

In two further aspects, the present invention features solid-state UV LDand UV LED devices, including at least two layers, each of which isindependently selected from AlN, GaN, InN, or binary or tertiary alloycombinations thereof. These devices are disposed atop a substrate of AlNhaving a dislocation density of about 10,000 cm⁻² or less, for example,5,000 cm⁻² or less, or, more particularly, about 1,000 cm⁻² or less. Incertain embodiments, the dislocation density is about 500 cm⁻², or aslow as 100 cm⁻². Indeed, as mentioned above, the invention facilitatesproduction of AlN substantially devoid of dislocation defects. In aparticular embodiment, the substrate has at least one regionsubstantially devoid of dislocation defects.

In yet another aspect, the invention is directed to a semiconductorstructure including an aluminum nitride single-crystal substrate havinga diameter greater than about 25 mm; and at least one layer epitaxiallygrown thereover. The epitaxial layer may comprise at least one of AlN,GaN, InN, or any binary or tertiary alloy combination thereof, whereinan average threading dislocation density at or near the surface of thesemiconductor heterostructure is less than about 10⁶ cm⁻² even when thedensity of misfit dislocations running parallel to the substrateinterface is much higher. The single-crystal substrate may have athickness of less than about 1 mm, and in some embodiments, of about 350μm. The single-crystal substrate preferably has an average dislocationdensity less than about 10,000 cm⁻². In some embodiments, at least onebuffer layer is disposed between the substrate and the strained layer.The buffer layer(s) may, for example, comprise AlN.

Active electronic devices may be fabricated on the semiconductorheterostructure. The device may, for example, be an optoelectronicdevice such as a laser diode (e.g., one having a maximum output at 405nm), or a high-brightness light-emitting diode. Alternatively, thedevice may be a non-optical electronic device such as a high power,radio-frequency amplifier. These devices can be expected to performbetter, have longer lifetime, and be produced with higher yield as aresult of having a lower threading dislocation density in the activedevice region. Preferably, an active electronic device in accordancewith the invention has a threading dislocation density of less than 10⁶cm⁻² in one or more device layers.

In still another aspect, the invention is directed toward a method offabricating a semiconductor heterostructure. The method comprises thesteps of providing an aluminum nitride substrate, growing a gradedbuffer layer thereover, and epitaxially depositing, over the gradedlayer, a layer comprising at least one of AlN, GaN, InN, or any binaryor tertiary alloy combination thereof, wherein the average threadingdislocation density near the surface of the semiconductorheterostructure or within the active device region of theheterostructure is less than about 10⁶ cm⁻² even when the density ofmisfit dislocations running parallel to the substrate interface is muchhigher. The substrate may be a single-crystal substrate having adiameter greater than about 25 mm and, in some embodiments, a thicknessof less than about 1 mm (e.g., about 350 μm). The single-crystalsubstrate preferably has an average dislocation density less than about10,000 cm⁻².

BRIEF DESCRIPTION OF THE DRAWINGS

In the drawings, like reference characters generally refer to the sameparts throughout the different views. Also, the drawings are notnecessarily to scale, emphasis instead generally being placed uponillustrating the principles of the invention. In the followingdescription, various embodiments of the present invention are describedwith reference to the following drawings, in which:

FIG. 1 is a graphical representation of crystal growth rate as afunction of nitrogen pressure, determined in accordance with variousaspects of the present invention;

FIG. 2 is a schematic elevational view of one embodiment of an apparatusof the present invention;

FIG. 3 is a schematic elevational view of an alternative embodiment of acomponent suitable for use with the apparatus of FIG. 2;

FIG. 4 is a schematic elevational view, on an enlarged scale, of thecomponent of FIG. 3;

FIG. 5 is a schematic elevational view of another embodiment of thecomponent of FIG. 3;

FIG. 6 is a cross-sectional view of a portion of a device fabricatedusing the method of the present invention.

FIG. 7 is a schematic elevational view of yet another embodiment of anapparatus of the present invention;

FIG. 8 is a schematic elevational view of a crucible for producing ahigh-purity source material in the source base crucible portion suitablefor use with the apparatus of FIG. 7;

FIG. 9 depicts an AlN single-crystal wafer fabricated in accordance withthe embodiments of the invention shown in FIGS. 7-8.

FIG. 10 is a schematic diagram of an apparatus for seeded bulk crystalgrowth of AlN, according to alternative embodiments of the invention.

FIG. 11A depicts a set of AlN single-crystal wafers fabricated inaccordance with the embodiments of the invention shown in FIG. 10.

FIG. 11B is a plot representing seed area expansions between 25-50% withrelatively small area seeds.

FIG. 11C depicts another AlN wafer fabricated in accordance with theembodiments of the invention shown in FIG. 10.

FIG. 12 is a schematic diagram depicting several device layersepitaxially grown on an AlN substrate fabricated in accordance withvarious embodiments of the invention.

DESCRIPTION OF THE INVENTION

As mentioned above, the present invention focuses on methods andapparatus for producing AlN substrates that advantageously have arelatively small lattice mismatch (around 2.2%) with respect to GaN, andhave an almost identical thermal expansion from room temperature to1000° C. Furthermore, AlN crystals grown in accordance with thetechniques disclosed herein have a low dislocation density of 10,000cm⁻² or less and may be grown to exceed 20-30 mm in diameter and 5 mm inlength. These AlN crystals also advantageously have the same wurtzitecrystal structure as GaN and nominally the same piezoelectric polarity.Also, the chemical compatibility with GaN is much better than that ofthe SiC. In addition, AlN substrates cut from the bulk crystals tend tobe attractive for Al_(x)Ga_(1-x)N devices requiring higher Alconcentration (e.g., for high temperature, high power,radiation-hardened, and UV wavelength applications).

Various aspects of the present invention stem from the recognition that,contrary to the teachings of the prior art, super-atmospheric pressuresmay be successfully utilized to produce single crystals of AlN atrelatively high growth rates and crystal quality. The inventiongenerally contemplates controlling one or more of (i) temperaturedifference between an AlN source material and growing crystal surface,(ii) distance between the source material and the growing crystalsurface, and (iii) ratio of N₂ to Al partial vapor pressures. Theinvention further contemplates increasing the N₂ pressure beyond thestoichiometric pressure to force the crystal to grow at a relativelyhigh rate due to the increased reaction rate at the interface betweenthe growing crystal and the vapor. This increase in the growth rate hasbeen shown to continue with increasing N₂ partial pressure untildiffusion of Al from the source to the growing crystal (i.e., thenegative effects of requiring the Al species to diffuse through the N₂gas) becomes the rate-limiting step. Various aspects of the presentinvention also contemplate a technique for establishing N₂ pressuressuitable for crystal growth, as will be discussed in greater detailhereinbelow. Moreover, employing higher-pressure nitrogen may have theadded benefit of reducing the partial pressure of aluminum inside thegrowth crucible, which may decrease corrosion within the furnace oftencaused by Al vapor inadvertently escaping the crucible.

In its various aspects, the present invention also features anapparatus, including a crystal growth enclosure or “crucible,” capableof providing a relatively sharp thermal profile, i.e., a relativelylarge thermal gradient along a relatively short axial distance, forsublimation and recondensation/nucleation. The crucible is configured tooperate at internal super-atmospheric pressures. During operation, thecrystal nucleates at a nucleation site, for example at the tip of thecrucible, and grows at a relatively high rate as the crucible movesrelative to the temperature gradient. As used herein, the term “axial”refers to a direction relative to the apparatus of the presentinvention, which is substantially parallel to push rod 17 shown in FIG.2. Also, the term “nucleation site” refers to a location of eitherseeded or unseeded crystal growth.

Before describing the apparatus and method of its operation in detail, adescription of the development and theory of the invention is in order.As early as 1998, the inventors developed a model of N₂ and Alincorporation into a growing AlN crystal. Particular aspects of thismodel are discussed in greater detail hereinbelow. It was found that theN₂ molecule is relatively difficult to incorporate into a growingcrystal, which leads to a small condensation coefficient for thismolecule. For this reason, the rate at which N₂ molecules collide withthe AlN surface (i.e., the N₂ flux) when the crystal is in equilibriumwith a mixed vapor of Al and N₂ is much greater than the rate at whichN₂ evaporates from the surface when the crystal is heated to the sametemperature in a vacuum (the so-called Langmuir evaporation rate). Usingonly the equilibrium pressures and the measured Langmuir evaporationrate, the model was able to correctly predict the maximum rate that AlNcrystals could be grown in atmospheric pressure N₂. More importantly,the model showed that the crystal growth rate, insublimation-recondensation growth if limited only by the surfacekinetics of N₂ incorporation, is directly proportional to the N₂pressure up to high pressures, which may be as high as 200 bar (20 MPa).This result is quite different from most crystal growth systems wherethe kinetically limited growth rate at the surface is maximized when thegas mixture above the growing crystal surface has the same stoichiometryas the growing crystal. It was also found that the Langmuir evaporationrate is almost 1000 times faster than the predicted AlN crystal growthrate at atmospheric pressure (even when diffusion is neglected). It isgenerally anticipated that one should be able to achieve growth ratescomparable to the Langmuir evaporation by going to stoichiometric gasmixtures. However, the AlN crystal growth system is very differentbecause of the difficulty in breaking the N₂ molecules into N atomswhich are then incorporated into the growing crystal.

While not wishing to be bound by a particular theory, the above resultcan be understood in view of the teachings of the present invention byrecognizing that the recondensation rate on the growing seed crystalgenerally matches the evaporation rate from the polycrystalline AlNstarting material during steady-state crystal growth. The AlN startingmaterial will sublime stoichiometrically and produce Al and N₂ vapor(the concentration of other gas species is believed to be too low toimpact the growth rate). Thus, by controlling the nitrogen pressureexternally, the Al partial pressure may rise until a steady state isachieved between evaporation at the hot surface and recondensation atthe cold surface. Because of the low condensation coefficient for N₂, ithas been found that the rate of evaporation/recondensation, as afunction of temperature, is greater at nitrogen pressures that exceedthe average stoichiometric pressure for AlN when the source material isAlN ceramic.

Referring to FIG. 2, at relatively low N₂ pressures, for realistictemperature differences between the polycrystalline starting material 11and the seed 7, this effect leads to relatively slow growth rates. Thisis a very different situation from crystal growth bysublimation/recondensation when both species have near unityaccommodation coefficients such as the case in SiC. Unfortunately, whenthe N₂ pressure exceeds the stoichiometric partial pressure for a givenpartial pressure of Al when stoichiometric, P_(N2)=0.5×P_(Al), masstransport of the Al, relative to the nitrogen gas, to the growingcrystal surface is generally needed. Thus, at some point, the growthrate becomes limited by diffusion of Al atoms through the gas phase eventhough the surface kinetics would continue to predict increased crystalgrowth rate with increasing N₂ pressure. Based on the inventors' presentunderstanding of the surface kinetics involved, it has been found thatthis cross-over point is only slightly greater than 1 atmosphere for thegrowth geometry that was used by Slack and co-workers in theaforementioned Slack reference, and as shown in FIG. 1. However, duringdevelopment of the present invention, it was found that this cross-overpoint is also dependent upon the diffusion lengths required for Altransport (which was approximately 2 to 5 cm in prior research). Byreducing this axial length in various embodiments of the presentinvention, which have been specifically configured to create a verysharp thermal profile in the work zone, it is has been found possible tosignificantly increase the growth rate relative to prior approaches.

Turning now back to FIG. 1, the predicted AlN growth rate is shown as afunction of N₂ pressure. The curves labeled with squares and crossesshow the growth rates assuming that it is limited by the diffusion of Al(for a 2.5 cm or 1 cm diffusion length, respectively) with noconvection, while the third curve shows the predicted growth rateassuming that the growth rate is limited by the surface kinetics ofnitrogen incorporation (ignoring gas-phase diffusion). The model assumesthat the AlN source material is at 2300° C. while the growing crystal ismaintained at 2200° C. These calculations also assume that the area ofthe evaporating surface and that of the growing crystal are equal andthat diffusion effects may be neglected. This last assumption, as shown,ceases to be true at high enough N₂ pressure. The cross-over pointgenerally depends on the experimental geometry.

Referring again to FIG. 2, many embodiments of the present inventionhave demonstrated that the diffusion issues described above can beaddressed, at least in part, by providing a net flow of gas from thesource 11 towards the growing crystal 7 greater than that caused simplyby the evaporation and recondensation process. This may be obtained whena thin-wall tungsten crucible 9 is used, and it may also be possible toobtain this effect with other crucible materials that are pervious tonitrogen gas, or with other types of selective barriers such as openings20, 21, described hereinbelow with respect to FIGS. 3-5. Nitrogen isable to diffuse through thin-walled tungsten crucibles at fairly highrates at the crystal growth temperatures (2300° C.) employed herein. Thediffusion rate of Al through the tungsten walls is much lower. Thus,under equilibrium conditions, the partial pressure of nitrogen insidethe crucible is identical to that outside the crucible, e.g., in chamber2 in FIG. 2, while the total pressure of gasses inside the crucible maybe higher due to the partial pressure of the Al vapor.

However, once crystal growth is initiated and the AlN source ismaintained at a higher temperature than the growing crystal, thenitrogen partial pressure at the cool end (e.g., location 19 of thegrowing crystal 7) of the crucible tends to become greater than at thehot end, while the opposite is true for the aluminum pressure.Meanwhile, the total gas pressure inside the crucible remainssubstantially uniform throughout to maintain mechanical equilibrium. (Asused herein, the term “mechanical equilibrium” refers to instances inwhich the total vapor pressures interior and exterior to the crystalgrowth chamber are substantially equal.) Thus, the nitrogen partialpressure at the cold end 19 of the crucible tends to exceed the nitrogenpressure outside the crucible (within chamber 2) while the opposite istrue at the hot end. For this reason, nitrogen tends to diffuse out ofthe crucible at the cold end and diffuse into the crucible at the hotend, resulting in a net flow of the gas mixture in the crucible from theAlN source toward the growing crystal.

While the understanding of this operation is new, its effect isconsistent with the experimental results achieved by Slack and McNellyand discussed in the Slack reference since their observed crystal growthrate of 0.3 mm/hr nominally could not have been achieved otherwise. Inaddition, Slack and McNelly observed that when pinholes were formed inthe CVD tungsten crucibles they used, the AlN vapor would escape. It hasbeen discovered by the present applicants that this occurs because, asexplained above, the total gas pressure in the crucible is greater thanthe nitrogen pressure outside the crucible. Once a pinhole opens up, thegas mixture in the crucible containing Al and N₂ starts being pushed outof the crucible. However, in this instance, nitrogen gas continues todiffuse from the outside through the walls of the crucible because thepartial pressure of N₂ is lower in the crucible than it was outside.Thus, the pressure in the crucible may be kept higher than the outsideeven when a pinhole is continuing to vent the gas mixture from thecrucible. The process typically continues until substantially the entireAl vapor inside the crucible is exhausted. Another source of additionalgas flow may be generated by the selective absorption of Al by thecrucible walls.

Thus, in particular embodiments of the present invention, a crucible maybe provided with one or more openings to facilitate nitrogen pressurecontrol/equilibration of the crucible, as will be discussed in greaterdetail hereinbelow with respect to FIGS. 3-5.

Growth Rate Model

As mentioned hereinabove, various aspects of the present invention stemfrom inventors' analysis of the applicable surface kinetics. The rate atwhich Al and N incorporate into the growing AlN has generally beenmodeled in two approaches. In the first, it is assumed that the nitrogenmolecules N₂ have a relatively low condensation coefficient compared tothe Al atoms due to a configurational barrier. In the second approach,it is assumed that the N₂ molecules physisorb onto the AlN surface.These N₂ molecules are then assumed to be kinetically hindered in theproduction of N atoms that can then incorporate into the AlN crystal.The inventors herein have modeled both approaches and both models leadto identical results.

The rate of change of the surface concentration [Al] of aluminum may begiven by:

$\begin{matrix}{\frac{\lbrack{Al}\rbrack}{t} = {{\beta_{Al}P_{Al}} - {C_{Al}\lbrack{Al}\rbrack} - {{B\left( {\lbrack{Al}\rbrack - \frac{K_{s}}{\lbrack N\rbrack}} \right)}.}}} & (1)\end{matrix}$

Likewise, the rate of change of the surface concentration [N] ofnitrogen atoms is:

$\begin{matrix}{\frac{\lbrack N\rbrack}{t} = {{2\; {\gamma\beta}_{N\; 2}P_{N\; 2}} - {C_{N}\lbrack N\rbrack}^{2} - {{B\left( {\lbrack{Al}\rbrack - \frac{K_{s}}{\lbrack N\rbrack}} \right)}.}}} & (2)\end{matrix}$

In these equations, the first term represents the addition of moleculesfrom the vapor. It is assumed that all of the Al atoms stick but only afraction y of the N₂ molecules condense on the surface. The term β_(i)represents the modified Hertz-Knudsen factor which is proportional tothe square root of mass of the i^(th) species (where i represents eitherAl or N₂) divided by the temperature. The condensation coefficient inthis model is not subscripted since we assume that it only applies tothe N₂ molecules.

In the equations above, the second term represents the evaporation ofmolecules into the vapor. The C_(i) term is a parameter introduced todescribe the rate of evaporation of the i^(th) species. Note that theevaporation of N₂ molecules proceeds like the square of the [N]concentration on the surface. Finally the last term represents theincorporation (or dissolution) of Al and N atoms into the crystal whereB is another parameter that is introduced. It is assumed that [Al] and[N] maintain an equilibrium concentration on the surface so as to obey

[Al]×[N]=K_(S)  (3)

where K_(S) is the equilibrium constant. It is also assumed that the Nand Al atoms on the surface come into equilibrium very rapidly, whichcan be modeled by making the assumption that B is very large. This tendsto constrain the [Al] and [N] concentrations to obey equation (3) almostexactly. The deviation from equilibrium tends to be just large enough sothat the time derivatives of [Al] and [N] are equal to zero once steadystate is achieved. The actual value of B is typically irrelevant underthese circumstances.

These rate equations lead to a relatively simple cubic equationdescribing the net flux F_(AlN) onto/off the surface of the AlN crystal,

$\begin{matrix}{{\left( {\frac{E_{L}^{2}P_{N\; 2}}{4\; \beta_{Al}^{2}P_{s}^{3}} - F_{AlN}} \right) \cdot \left( {{\beta_{Al}P_{Al}} - F_{AlN}} \right)^{2}} = {E_{L}^{3}.}} & (4)\end{matrix}$

Significantly, the only parameters appearing in this cubic equation arethe Langmuir evaporation rate E_(L) and the stoichiometric nitrogenpressure P. It is rather remarkable that, in the limit of very large B,of the five free parameters that appear in Eqs. (1) and (2), only twoparameters are needed in Eq. (4) to determine the net flux at anynitrogen and aluminum partial pressure. Both of these parameters havebeen determined experimentally although there is substantial uncertaintyin the Langmuir-evaporation rate as pointed out earlier.

These equations thus support the use of increasing nitrogen pressurebeyond stoichiometric pressures, which is contrary to prior teachings asdiscussed hereinabove. In the limit of a small value of y, on the orderof 10⁻⁵, it is possible to linearize this expression to give F_(AlN) asa linear function of P_(N2). The growth rate can be determined bysetting

F _(AlN)(T _(H))=−F _(AlN)(T _(C)).  (5)

A simple expression for this growth rate that depends on the differencein temperature Δ_(T)=T_(H)−T_(C) and on the nitrogen pressure, which maybe controlled independently, has been determined. If the effects ofdiffusion are ignored (so that the partial pressures of N₂ and Al areconstant in the region between the source and seed), for a given Δ_(T)and N₂ pressure, the Al pressure tends to adjust itself to some valueintermediary to the equilibrium Al pressures at the hot and coldsurfaces so as to force Eq. (5) to be true. In this case,

F _(AlN) =A _(o)(ΔT,T _(H))P _(N2).  (6)

The value of F_(AlN) is linear in Δ_(T) for Δ_(T)<100° C. It has beenfound that the linear dependence of F_(AlN) on P_(N2) is valid (within1%) between 0 and 200 bars of nitrogen gas for the experimentallydetermined values of E_(L) and P, in the temperature range 2000°C.<T_(H)<2500° C. It has been determined that

A _(o)(ΔT=50° C.,T _(H)−2300° C.)=0.156 mm−hr⁻¹−bar⁻¹,  (7)

using the data of Bolgar et al. (A. S. Bolgar, S. P. Gordienko, E. A.Ryklis, and V. V. Fesenko, in “Khim. Fiz. Nitridov”, p. 151-6 (1968)[Chem. Abstr. 71, 34003] (1969)]) for the Langmuir evaporation rate (259mm/hr) and the equilibrium, stoichiometric nitrogen pressure determinedfrom the JANAF tables (M. W. Chase et al. “JANAF Thermochemical Tables”,Third Edition (1985), see J. Phys. Chem. Ref. Data 14, Supplement No. 1(1985)), P_(S)(2300° C.)=0.13 bar. This equilibrium pressure would leadto an effective growth rate of 18.9 m/hr if all the nitrogen andaluminum stuck to the surface without re-evaporation andP_(Al)=2P_(N2)=P_(s), such as shown in FIG. 1. These calculations assumethat the area of the evaporating surface and that of the growing crystalare equal and that diffusion effects may be neglected. This lastassumption is rather important and, as FIG. 1 shows, ceases to be trueat high enough N₂ pressure.

The observed growth rate in the Slack reference for T_(H)=2300° C. andT_(C)=2200° C. run in 0.95 bar of N₂ plus 0.05 bar of H₂ was 0.30 mm/hr.This should be compared with the theoretically determined growth rate of0.32 mm/hr. This is a remarkable agreement given the uncertainties inthe experimental data for the Langmuir evaporation rate and in themeasured growth rate by Slack and McNelly. There are no adjustableparameters in the way that this theory was developed, the theory onlydepends on the experimentally-determined equilibrium pressures and themeasured Langmuir evaporation rates. Notably, the experiment wasconducted in a crucible where the growing crystal surface was smallerthan the evaporating AlN source material. This tends to lead to anamplification of the observed growth rate. It should also be noted thatthese equations predict a theoretical growth rate of 0.020 mm/hr at thestoichiometric nitrogen pressure (0.13 bar) at a source temperature of2300° C. and a Δ_(T) of 50° C.

Crystal Growth Furnace

Turning now to FIG. 2, the apparatus of the present invention, forexample, a furnace, includes a heating source 6, such as aradio-frequency (“RF”) coil for inducing an electro-magnetic fieldwithin a growth chamber 2. This electro-magnetic field couples with ametallic susceptor, for example a push tube 3 located concentricallyinside the coil and provokes heat generation thereon by the Jouleeffect. Although in a particular embodiment, susceptor/tube 3 iscylindrical, i.e., has a circular axial cross-section, as used herein,the terms “tube” or “tubular” also include tubes of non-circular axialcross-section. The relative position and dimension of the push tube withrespect to the shielding elements and coil creates a thermal gradientalong the walls of the susceptor 3, i.e., in the axial direction. Acrucible 9 is movably disposed concentrically within tube 3, andincludes the high-purity source material 11 at a proximal end thereof,e.g. polycrystalline AlN, and eventually, the growing AlN crystal 7 atthe distal end (e.g., at tip 19) thereof.

The crucible 9 may be fabricated from material such as tungsten orrhenium or tantalum nitride, having walls such as discussed hereinabove,which are thin enough to form a barrier which selectively permits andprevents the respective diffusion of nitrogen and aluminum,therethrough. Alternatively, the crucible may be configured withopenings configured to effect similar selectivity, as discussed ingreater detail hereinbelow. In some embodiments, metallic baffles 13, 4are disposed at proximal (e.g., bottom) and distal (e.g., top) ends,respectively, to facilitate control of the temperature gradient alongthe crucible longitudinal axis. One may selectively remove some of thebaffles from the distal end relative to the proximal end to achieve thedesired temperature gradient. The top set of baffles may include acenter hole that facilitates temperature measurement at the tip of thecrucible using a conventional pyrometer 1. If the operational conditionsare adequate in terms of pressure, temperature, and displacement rate ofthe crucible respect to the gradient, a single crystal boule 7 is formedat the distal end of crucible 9, i.e. tip 19.

In various embodiments, this invention focuses on the arrangement of theshielding elements in the system that facilitates setting a desirablethermal gradient along the walls of push tube 3. As shown, in variousembodiments, there are two distinct sets of shielding elements aroundthe push tube. In a particular embodiment, the first set includes twoconcentric open joint tungsten cylinders 8, each having a thickness ofless than about 0.005″. Skilled artisans will readily recognize thatother refractory metals such as molybdenum or rhenium may be substitutedfor tungsten.

As used herein, the term “open joint” refers to the cylinders 8 havingan open longitudinal seam (i.e., the cylinders to not extend a full 360degrees) so that there is no continuous electrical path around thecylinder. This serves to prevent the cylinders from coupling to the RFfields and becoming too hot and/or absorbing power intended for heatingthe crucible. Both the open joint and the thickness are chosen tominimize the coupling between these metallic parts and theelectromagnetic field. In addition, in various embodiments, cylinders 8are preferably axially shorter than RF coil 6 and located nominally inthe center of the coil (e.g., both concentrically with the coil andequidistantly from both axial ends thereof) to avoid inducing localnon-axial-symmetric hot spots on the shields. The second set of shields10 preferably includes pyrolytic boron nitride (pBN) cylinders (e.g., incertain embodiments, between five and seven), which are preferably atleast 0.050 inches (1.3 mm) thick and are several centimeters longerthan the RF coil. While one of the purposes of the pBN shields 10 is tothermally insulate the tube 3 to obtain the desire temperature in thework zone, the mission of the metallic shields is two-fold. Shields 8serve as heat reflectors causing the temperature in the center of thehot zone to be much higher than at the coil ends. In addition, theyserve to protect the push tube 3 from picking up boron generated by thecontinuous sublimation of the pBN shields 10. (Boron has a eutecticpoint with tungsten at the growth temperatures used in this invention ifthe boron concentration exceeds a certain value. Once the boron pickedup by the push tube is higher than that value, a liquid phase tends toform on the skin of the push tube leading to its failure.) In variousembodiments of the invention, the shielding arrangement describedhereinabove advantageously produces a sharp thermal gradient on thecrucible 9 (e.g., over 100° C./cm) as its tip 19 moves axially beyondthe metallic shields 8. As discussed hereinabove, this relatively largegradient has been shown to facilitate large growth rates. As also shown,the shields 8, 10 may be enclosed in a dielectric tube 5, fabricatedfrom quartz or other suitable material, to insulate the coil 6 from themetallic elements of the system and prevent arcing or other undesirableinteractions with them.

Still referring to FIG. 2, in many embodiments of the invention, thefurnace includes a crucible pedestal 12 disposed within push tube 3, forsupportably engaging crucible 9 within tube 3. The push tube is itselfsupported by a push tube pedestal 14, which is engaged by a push rod 17for axial actuation. Shielding pedestal 15 supportably engages pBNshields 10, and gas outlets and inlets 16 and 18, respectively enablevacuum/venting and pressurization.

In some embodiments of the present invention, the apparatus may alsoinclude a controller 24, such as one or more computers ormicroprocessors and related software, coupled to actuators and the like,suitable to automate at least some of the aspects the operation of thefurnace. For example, as shown, controller 24 may be communicablycoupled to heater 6 and to pyrometer 1 forming a conventionalclosed-loop system capable of maintaining chamber 5 at desiredtemperatures in response to execution of various temperature ramproutines by the controller. Controller 24 may be similarly coupled to aconventional vacuum pump 28, and to an electronically-activatable valve30 disposed at gas outlet 16 to automatically effect evacuation andventing of chamber 2. The controller may also be coupled to valve 32disposed at gas inlet 18 to effect pressurization of the chamber, whilethe controller is also coupled to an actuator 34 to automaticallycontrol operation of push rod 17. Skilled artisans will readilyrecognize that controller 24 may also be coupled to conventionalpressure and position sensors (not shown) disposed within the chamber toprovide closed-loop feedback to control the gas pressure and push rodposition.

In still further embodiments, crucible 9 may be provided with one ormore openings 20, 21 that allow the nitrogen partial pressure inside thecrucible to be controlled by the nitrogen pressure outside the crucible,while minimizing the amount of Al vapor that escapes from the crucible.This goal may be accomplished by allowing the nitrogen to diffusethrough a relatively thin-walled tungsten crucible, such as the onedisclosed in the Slack reference cited above. Such a crucible may bewelded shut so the only route for Al to escape the crucible is bydiffusion through the walls. Since N₂ is able to diffuse through W muchmore rapidly than Al, this procedure was effective. In this event, thewalls of the crucible need to be kept relatively thin to allow the N₂inside the crucible to equilibrate with the N₂ outside in a reasonableamount of time. Otherwise, the time required to equilibrate (i.e.,through relatively thick-walled crucibles) may be prohibitive.Disadvantageously, however, relatively thin-walled tungsten cruciblestend to have a lifetime that is too short to grow relatively largecrystals. Other disadvantages of this thin-walled approach include theneed to evacuate the crucible prior to electron-beam welding, and theexpense of such welding.

Thus, embodiments of the present invention contemplate mechanicalapproaches to allow the nitrogen to equilibrate, while minimizing theescape of Al vapor. Turning now to FIGS. 3-4, in alternativeembodiments, a two-piece crucible 9′, which includes a base portion 25,which is threadably engageable with a tip portion 27, is employed. Thisapproach includes providing (e.g., by drilling) a hole 20 in crucible 9′to allow nitrogen flow during heating phases 40, 44 (describedhereinbelow), but which is small enough to allow only minimal diffusionof Al vapor through the hole once the total pressure in the crucible isequal to the total pressure outside the crucible. As best shown in FIG.4, in a particular exemplary embodiment, this may be accomplished bydrilling a 25-mil (0.63 mm) diameter hole 20 into proximal end 25 ofcrucible 9′, (although it may also be provided elsewhere, such as indistal end 27) and then using a 20 mil (0.50 mm) diameter wire 23 toplug the hole. The nominally 0.13 mm clearance (best shown in FIG. 4)provided by this assembly has been shown to successfully allow thenitrogen pressure to equilibrate prior to growth and, during growth.Moreover, it is small enough so that the diffusion of Al vapor throughit is acceptably low, or so small that it effectively becomes pluggedduring growth, to substantially prevent Al diffusion. Skilled artisanswill readily recognize that this crucible 9′ may be used with or withouthole 20 and wire 23.

Turning now to FIG. 5, as mentioned above, another approach is toprovide a horizontal-seal (opening) 21, by providing a two-piececrucible 9″ including the proximal end 25′ containing source 11, and thedistal end 27′ containing crystal 7. The mating surfaces of the two endsthat form seal 21 are configured to be sufficiently smooth so that thereis little, if any, ability for gas to diffuse therethrough if theinterior and exterior crucible pressures are at equilibrium. However,the seal 21 will permit (i.e., be unable to resist) gas flow from eitherinside the crucible or outside the crucible if the pressures are notsubstantially equal. As shown, the seal is disposed in a substantiallycentral region of the crucible where no deposition of AlN will occur,e.g., by disposing the seal in the hottest region of the crucible 9′during crystal growth. This disposition of the seal effectively preventsthe AlN growth from opening seal 21 and thereby permitting unacceptablyhigh levels of Al vapor diffusion therethrough.

Having described the principles and apparatus of the present invention,the method of operation, i.e., the actual growth process using thesystem described above is now described in conjunction with thefollowing Table 1.

TABLE 1 30 Evacuate chamber 32 Refill chamber with nitrogen gas 34Repeat steps 30, 32 36 Pressurize chamber to I bar with a gas comprisingabout 95-100% N₂ and 0-5% H₂ 37 Place source material in proximal end ofcrucible 38 Place crucible in tube, with its tip in the high-temperature region at proximal end of the shield 40 Increase temperatureof the crucible to about 1800° C. 42 Maintain gas at a predeterminedsuper-atmospheric pressure 44 Ram temperature to growth temperature 46During step 44, continuously adjust pressure to maintain it at thepressure of step 42 48 Once growth temperature is reached, pushrod isactuated to move crucible tip towards distal end of chamber 50 Maintainconstant pressure during step 48 52 Stop movement of the push rod 54Cool furnace to room temperature

As recited in Table 1 above, in some embodiments, crystal growthinitially involves evacuating the chamber 2 (FIG. 2) (step 30), e.g., topressures on the order of about 0.01 mbar (1 Pa) using a vacuum pump.The chamber is then refilled with nitrogen gas (step 32). This step ispreferably repeated several times to minimize oxygen and moisturecontamination (step 34). The chamber is then pressurized to about 1 bar(100 kPa) with nitrogen gas which is preferably mixed with a smallamount of hydrogen (step 36). For example, a gas including about 95-100%N₂ and 0-5% H₂ is suitable. In particular embodiments, acommercially-available mixture of about 3% H₂ and 97% N₂ is employed.Polycrystalline AlN source material is placed at a proximal end of thecrucible (step 37). The crucible 9 may then be evacuated and sealed, ormay be provided with openings as described hereinabove. The crucible isthen disposed concentrically within the tube 3, with tip 19 in thehigh-temperature region of the furnace (i.e., nominally within theproximal end of shield 8) (step 38). The temperature is then increasedto bring the tip of the crucible to a temperature of approximately 1800°C., in particular embodiments, within about 15 minutes (step 40). At theend of this temperature ramp, the gas pressure is set and maintained ata predetermined super-atmospheric pressure (step 42), and thetemperature is ramped to a final crystal growth temperature (step 44),e.g., in about 5 hours. During this ramping, the pressure iscontinuously adjusted, e.g., using a vent valve (not shown) to maintainit at that fixed value (step 46). The goal of this ramping is to enhancethe purity of the source material by permitting part of the oxygen stillcontained within it to diffuse out of the crucible (e.g., through thecrucible walls). This diffusion occurs because the vapor pressure of thealuminum suboxides (such as Al₂O, AlO, etc.), generated due to thepresence of oxygen in the source material, and is known to be higherthan that of Al over AlN for the same temperature.

Once the growth temperature is reached, the push rod is actuated to movethe crucible tip towards the distal end of the chamber, and relative tothe thermal gradient (step 48). As discussed hereinabove, the crucibletip is initially located within the highest temperature region of thesusceptor 3 at the beginning of the growth run. As the push tube movesupwards (i.e., towards the distal end of the furnace 2, tip 19 becomescooler than the source material which promotes effective mass transportfrom the source material to the colder tip of the crucible. As shown inFIG. 2 and described above, the push tube, including crucible 9 disposedtherein, is moved axially upon actuation of push rod 17. However, pushrod 17 may alternatively be configured to move the crucible axiallyrelative to the tube, without departing from the spirit and scope of thepresent invention.

During the growth process, the pressure is maintained at a constantpredetermined value (step 50). The most appropriate value for thispressure typically depends on the axial spacing between the sourcematerial 11 and the (closest) surface of the growing crystal 7, as wellas the rate of nitrogen diffusion through the crucible walls or flowthrough the predrilled holes. It may also be appropriate to activelyadjust the gas pressure over a relatively narrow range during crystalgrowth to compensate for any changes in the spacing between the surfaceof the sublimating source and growing crystal surface.

In particular embodiments, a pressure of about 18 psi has been used todemonstrate growth rates of 0.9 mm/hr with the source material 11 tocrystal surface 7 separation of approximately 2 cm, employing tungstencrucibles fabricated by either chemical vapor deposition or powdermetallurgy technique (such as those described in commonly assigned U.S.Pat. No. 6,719,843 entitled “Power Metallurgy Tungsten Crucible for AlNCrystal Growth” which is fully incorporated by reference herein). Thesource to growing crystal surface distance may vary during the growthrun if the area of the growing crystal surface is different than thesurface area of the source material and the growth rate (i.e., axialrate of movement of the crucible through the temperature gradient) mayneed to be adjusted to account for this change. However, typically thesurface area of the source and growing crystal surface will be keptnominally constant and approximately the same size so that theseparation between the source and growing crystal surface will remainabout constant during most of the growth.

Finally, the movement is stopped (step 52) and a cooling ramp (step 54)is established to bring the system to room temperature. Using pressuresin the range 100 kPa to 150 kPa (1 atm to 1.5 atm); single crystalboules have been grown at an axial pushing rate ranging between about0.4 and 0.9 mm/h, for example, at the rate of 0.455 mm/h. By adjustingthe distance between the source material and the growing crystalsurface, and by adjusting the temperature gradient, other useful growthconditions may be obtained. Hence, skilled practitioners may usefullyuse various embodiments of the present invention with total chamberpressures from 50 kPa to 1 MPa (0.5 atm to 10 atm) and axialpushing/growth rates of 0.3 to about 3 mm/h.

Advantageously, various embodiments of the present invention feature anAlN crystal growth system capable of controlling the N₂ partial pressureindependently of the Al partial pressure, while substantially preventingsignificant amounts of Al from escaping from the growth crucible. Theseembodiments thus teach and facilitate use of partial nitrogen pressuresgreater than stoichiometric, and the use of total vapor pressures atsuper atmospheric levels (i.e., greater than one atmosphere).

Furthermore, by slicing or cutting the bulk AlN crystals of the presentinvention, crystalline AlN substrates of desired thickness—for example,about 500 μm or 350 μm—can be produced. These substrates can then beprepared, typically by polishing, for high-quality epitaxial growth ofappropriate layers of AlN, GaN, InN and/or their binary and tertiaryalloys to form UV LDs and high-efficiency UV LEDs. The aforementionednitride layers can be described by the chemical formulaAl_(x)Ga_(y)In_(1-x-y)N, where 0≦x≦1 and 0≦y≦1-x.

The surface preparation of the AlN is important to obtain high-qualityepitaxial growth of the nitride layers on the AlN substrate. Surfacedamage is preferably carefully removed in order to obtain high-qualityepitaxial layers needed for fabrication of high performance nitridesemiconductor devices. One successful approach to remove surface damagefrom the AlN substrate is to employ a chemical-mechanical polishing(CMP) approach, e.g. as described in U.S. Pat. No. 7,037,838 (the “838patent”), incorporated herein by reference. Through this approach, itpossible to produce very high-quality epitaxial layers ofAl_(x)Ga_(y)In_(1-x-y)N using organo-metallic vapor phase epitaxy(OMVPE) with low dislocation densities, particularly when x exceeds 0.5.Those skilled in the art will recognize that other epitaxial growthtechniques such as molecular beam epitaxy (MBE) or hydride vapor phaseepitaxy (HVPE) may also be successfully employed to produce high-qualityepitaxial layers on the low dislocation density AlN substrates of thepresent invention.

The present invention also features a method for making solid-state UVlaser diodes and UV light emitting diodes, as well as the devicesthemselves. The LDs and LEDs are fabricated on AlN substrates made bythe present process, preferably emitting at wavelengths ranging fromabout 200 nm to about 320 nm. Furthermore, the solid state UV LEDsfabricated via some of the embodiments of the claimed method exhibithigh efficiency wherein the fraction of electric power converted into UVradiation (total energy per unit time delivered to the device) and thenconverted into UV radiation power (total energy radiated per unit timein the form of UV photons) exceeds 10%.

The process for fabricating these devices in compound semiconductors mayinclude the following steps. Typically, a p-n junction is formed bygrowing p-type and then n-type (or vice-versa) epitaxial layers on anappropriate substrate. Metal contacts (preferably, ohmic contacts) areattached to the p-type and n-type semiconductor layers. The LED or LDfunctions by passing current through the device in the “forwarddirection.” That is, a power source pushes electrons out of the n-typesemiconductor toward the p-type semiconductor and holes out of thep-type semiconductor toward the n-type semiconductor. The electrons andholes recombine with each other to produce photons at a wavelength thatis determined by the bandgap energy of the semiconductor region wherethe recombination occurs and may be shifted to somewhat higher energiesby quantum confinement effects as well as by strain and impurities inthe recombination region to higher or lower energies. Many factors mayaffect the efficiency of the LEDs. However, a very significant factor isthe efficiency with which electrons and holes, once either generated orpushed out of their equilibrium positions, recombine to produce thedesired radiation. In addition, LDs require an optical cavity to allowamplification by stimulated emission of the appropriate photons. Thisefficiency is improved by defining the region of recombination, often bycreating one or more quantum wells. However, defects, such asdislocations, provide non-radiative recombination sites and may reducethe recombination efficiency. Most importantly, once the density ofnon-radiative recombination centers exceeds the typical diffusion lengthof carriers before they recombine, the loss in efficiency will becomevery significant and, as the density of defects is increased evenfurther, the device may either perform badly, i.e. very inefficiently,or not at all.

As skilled artisans will readily recognize, the addition of In to III-NLED and LD devices facilitates localization of the free carriers awayfrom non-radiative defects in recombination region probably through theformation of composition fluctuations where regions with higher Incontent have a somewhat smaller bandgap (and hence, are a lower energyregion for the free carriers to localize). This enables visible LEDs(emitting in the wavelength range from 400 to 550 nm) to function eventhough the dislocation density is very high (exceeding 10⁸ cm⁻²) (see,for example, J. Y. Tsao, Circuits and Devices Magazine, IEEE, 20, 28(2004)). These high dislocation densities are the result of growingnitride epitaxial layers on foreign substrates such as sapphire(crystalline Al₂O₃) or SiC. The wavelength of emission from pure GaN isapproximately 380 nm and decreases to 200 nm for pure AlN. Wavelengthsin between these two end points can be achieved by alloying AlN withGaN. InN can also be introduced to make an Al_(x)Ga_(y)In_(1-x-y)Nalloy. While success has been achieved in making LEDs at wavelengthsdown to 250 nm, the efficiencies of these devices remains quite poor(<1%) (A. Khan, “AlGaN based Deep Ultraviolet Light Emitting Diodes withEmission from 250-280 nm,” presented at the International Workshop onNitride Semiconductors in Pittsburgh, Pa. (Jul. 19, 2004)).

Furthermore, low levels of dislocation densities in the substrate appearto be critical to the fabrication of high efficiency UV LDs and LEDs,particularly at wavelengths shorter than 340 nm (S. Tomiya, et al.,“Dislocations in GaN-Based Laser Diodes on Epitaxial Lateral OvergrownGaN Layers,” Phys. Stat. Sol. (a) 188 (1), 69-72 (2001); A. Chitnis etal., “Milliwatt power AlGaN quantum well deep ultraviolet light emittingdiodes,” Phys. Stat. Sol. (a) 200 (1), 88-101 (2003); M. Takeya, et al.,“Degradation in AlGaInN lasers,” Phys. Stat. Sol. (c) 0 (7), 2292-95(2003), all incorporated herein by reference. Nitride semiconductorlayers grown on sapphire substrates typically have dislocation densitiesof about 10⁸-10¹⁰ cm⁻², and, recently, S. Tomiya et al. have producedELO GaN substrates with dislocation densities of about 10⁶ cm⁻², asdiscussed in the aforementioned journal article. Furthermore,high-efficiency UV LEDs and UV LDs require layers ofAl_(x)Ga_(y)In_(1-x-y)N with relatively high aluminum concentration.

Higher efficiency LEDs have been recently fabricated on the AlNsubstrates prepared according to the embodiments of the presentinvention, exhibiting an improvement over an order of magnitude inphotoluminescence efficiency of the low-defect layers grown on thesubstrates cut from bulk crystals fabricated in accordance with variousembodiments of the present invention compared to the high-defect densityepitaxial layers grown on sapphire (which has a defect density greaterthan 1,000,000 dislocations cm⁻²). The improved results are believed tobe due to the low levels of dislocations densities in substrates cut andprepared from bulk AlN crystals grown in accordance with variousembodiments of the present invention. Specifically, as mentioned above,the present invention contemplates the fabrication of single crystals ofAlN having dislocation defect densities of about 10,000 cm⁻² or less. Insome embodiments, the dislocation densities are about 5,000 cm⁻² orless, or, more preferably, about 1,000 cm⁻² or less on average. Inparticular versions of these embodiments, the dislocation densities inat least some of the areas of the AlN crystal are about 500 cm⁻² orless, for example, as low as 100 cm⁻², with certain regions—for example,exceeding 5×5 mm²—being substantially devoid of dislocation defects,i.e. having no visible dislocations at all. This low dislocation densitycan be measured by Synchrotron White Beam X-ray Topography (SWBXT) (SeeB. Raghothamachar et al., J. Crystal Growth 250, 244-250 (2003); and L.J. Schowalter et al., Phys. Stat. Sol. (c) 1-4 (2003)/DOI10.1002/pssc.200303462), both articles incorporated herein by reference.

FIG. 6 and the following example illustrate the fabrication of anexemplary UV LD emitting at a wavelength of 260 nm using conventionaltechniques. The LD is disposed on a low dislocation density AlNsubstrate formed according to some of the embodiments of the presentinvention. The orientation of the AlN substrate in the example isc-face, Al polarity. However, as one of ordinary skill will readilyrecognize, other orientations, such as N polarity, m-face, and a-face,could be used instead, depending on the desired application.Furthermore, as one of ordinary skill will also recognize, many other LDand LED devices emitting at various other wavelengths can alternativelybe prepared by the method disclosed herein. In addition, thesubstitution of multiple quantum wells in the active region as well aspatterning the surface to laterally confine the radiation in the quantumwells are well-known procedures to enhance the performance of the laserdiode. The example is included to illustrate the fabrication of one suchdevice, and the invention is not limited to the device illustrated inthe example.

Example—260 nm Laser Diode Fabrication

FIG. 6 is a schematic cross-sectional view of a portion of a 260-nmlaser diode fabricated using the method of the present invention.Initially, low defect density AlN substrate 40, which is prepared usingthe method discussed above, is polished by CMP. The polished substrateis then introduced into a conventional OMVPE system. The surface of thelow defect density AlN substrate is cleaned to remove any oxide or otherimpurities on the crystal surface by heating the substrate at atemperature of 1150° C. for 20 min under ammonia plus nitrogen orhydrogen plus nitrogen flow prior to growth. An epitaxial layer 42 ofAlN having a thickness of about 100 nm is then grown on the substrate toimprove its surface quality before starting to grow the device layers.Next, an undoped Al_(x)Ga_(1-x)N buffer layer 44 having a thickness ofapproximately 1 μm is grown atop the epitaxial AlN layer to relievelattice mismatch through the formation of misfit dislocations. Formationof threading dislocations, which will continue to propagate through thedevice layers, is minimized by grading x from 1 to the final value of0.5 (i.e. to 50% Al concentration). Onto the buffer layer, a 1-μm thicklayer 46 of Si-doped Al_(x)Ga_(1-x)N (x=0.5) is then grown to providethe n-type contact to the LD. A 50-nm thick layer 48 of Si-dopedAl_(x)Ga_(1-x)N (x=0.6) is then grown onto the Si-doped Al_(x)Ga_(1-x)N(x=0.5) layer 46, followed by the growth of a 10-nm thick layer 50 ofundoped Al_(x)Ga_(1-x)N (x=0.5). Then, a 50-nm thick layer 52 ofMg-doped Al_(x)Ga_(1-x)N (x=0.6) is grown onto the layer 50 followed bythe growth of a 1-μm thick layer 54 of Mg-doped Al_(x)Ga_(1-x)N (x=0.5).After the growth steps, the substrate 40 (now having epitaxial layers42, 44, 46, 48, 50, 52 and 54 thereon) is slowly ramped down from thegrowth temperature of about 1100° C. and removed from the OMVPE system.In some embodiments, the top surface of epitaxial layer 54 is thencoated with a metal contact layer 56 for the p-type semiconductor, andmetal layer is coated with a photoresist (not shown), which is thendeveloped. The photoresist is removed where the n-type metal contact 58will be formed. The substrate, along with the epitaxial layers and themetal layer thereon, are then etched such that the semiconductor isremoved down to the n-type layer 46, which will be used for the n-typemetal contact 58. A second coating of photoresist (for lift off) (notshown) is deposited, which is then patterned and removed where then-type contacts are desired. The n-type metallization is now complete,and the metal coating adjacent the second photoresist layer is removedto produce the desired wafer. Laser facets are achieved by cleaving thewafer. These facets may optionally be coated to protect them fromdamage. Wire bonding contacts (not shown) are made to the p-type andn-type metal layers and the laser diode is packaged appropriately. Inother embodiments, the mesa etching and n-contact fabrication precedesp-type contact metallization due to the higher annealing temperature forthe n-type contact (−900° C.) compared to the p-contact anneal at ˜600°C.

The growth of the bulk single crystal of AlN has been described hereinprimarily as being implemented by what is commonly referred to as a“sublimation” technique wherein the source vapor is produced at least inpart when crystalline solids of AlN or other solids or liquidscontaining AlN, Al or N sublime preferentially. However, as alsodisclosed herein, the source vapor may be achieved in whole or in partby the injection of source gases or like techniques that some wouldrefer to as “high-temperature CVD”. Also, other terms are sometimes usedto describe these and other techniques that are used to grow bulk singleAlN crystals according to this invention. Therefore, the terms“depositing,” “depositing vapor species,” and like terms will sometimesbe used herein to generally cover those techniques by which the crystalmay be grown pursuant to this invention.

Thus, the AlN boules fabricated using the embodiments describedhereinabove may be used to produce substrate by cutting a wafer orcylinder from the bulk single-crystal AlN, preparing a surface on thewafer or cylinder in a known manner to be receptive to an epitaxiallayer, and depositing an epitaxial layer on the surface usingconventional deposition techniques.

In particular embodiments of the invention, large, e.g. greater thanabout 25 mm in diameter, single-crystal AlN wafers are produced fromsingle-crystal AlN boules having a diameter exceeding the diameter ofthe final substrate, e.g. boules having a diameter greater than about 30mm. Using this approach, after growing the boule and orienting it, e.g.by employing x-ray Laue diffraction technique, to obtain a desirablecrystallographic orientation for the wafer, the boule is mechanicallyground down to a cylinder having a desirable diameter and then slicedinto individual wafers, e.g., using a wire saw. In some versions ofthese embodiments, the boules can be grown by, first, producinghigh-quality single-crystal seeds, and then using the seed crystals asnuclei to grow larger diameter single-crystal boules through a crystalexpansion growth run. It is then possible to take large-diameter slicesfrom this second crystal growth process and use them to growlarge-diameter crystals without diameter expansion. In alternativeversions, the crystal growth is self-seeded, i.e. the crystal is grownwithout employing single-crystal seeds. Various versions of the aboveembodiments are described in more detail below.

Self-Seeded Crystal Growth

Referring to FIGS. 7-8, in some embodiments, an apparatus 70 forself-seeded growth of single-crystal AlN boules includes a crucible 72having a conical crucible portion 74, such as the one disclosed in U.S.Pat. No. 6,719,842 (“the '842 patent”), incorporated herein byreference, consisting essentially of tungsten and fabricated by a CVDprocess. Multiple grain layers within the wall of the conical portioncan be obtained by interrupting the tungsten deposition several timesbefore the final wall thickness was obtained. Other crucible materialscan be used such as tungsten-rhenium (W—Re) alloys; rhenium (Re);tantalum monocarbide (TaC); a mixture of Ta₂C and TaC; a mixture ofTa₂C, TaC and Ta; tantalum nitride (Ta₂N); a mixture of Ta and Ta₂N;hafnium nitride (HfN); a mixture of Hf and HfN; a mixture of tungstenand tantalum (W—Ta); tungsten (W); and combinations thereof. In certainversions of these embodiments, a tip 75 of the conical portion has anarrow angle, for example, about 15°. The apparatus further includes asource base crucible portion 76, having an AlN source 77, for example,consisting essentially of high-purity AlN ceramic disposed therein. Invarious embodiments, the source base crucible is fabricated fromtungsten and is prepared from a high-density, powder metallurgy cylinderhollowed out by electric discharge machining. Thus fabricated, the basecrucible portion includes a plurality of grains arranged in multiplelayers within the wall of the crucible.

Referring to FIG. 8, the high-purity source material in the source basecrucible portion 76 can be produced in a crucible 80 by reactinghigh-purity Al (e.g. having 99.999% purity, available from Alpha Aesarof Ward Hill, Mass., USA) with high-purity N₂ gas (e.g. having 99.999%purity, available from Awesco of Albany, N.Y., USA). In a particularversion, chunks of high-purity AlN ceramic 82, e.g. weighing about 9 gor less, are placed in a bottom portion 84 of the crucible 80 and heatedto about 2300° C. in a forming gas atmosphere in order to sublime theAlN and recondense it in the source base portion 76. As a result, thedensity of the resulting ceramic can be increased to approximatelytheoretical density by sublimation transport to decrease the surfacearea relative to the volume of the source material. The resulting AlNceramic material may have impurity concentration of less than about 500ppm.

Referring again to FIG. 7, in various embodiments, the source base andconical portions are then assembled together, e.g. employing ahorizontal seal, forming the crucible 72. The mating surfaces of thesource base and the conical portion are preferably treated prior toassembly by polishing with a diamond grit (e.g. by a 15 micron gritfollowed by 1 micron grit) on a polishing wheel and then by 1200 gritSiC sand paper until extremely smooth to improve quality of the joint.

The crucible loaded with the source material can then be assembled inthe furnace, e.g. high-pressure crystal growth furnace available fromArthur D. Little, Inc. In a manner analogous to the embodiments of theapparatus described above in connection with FIGS. 2-5, the crucible canbe placed on a pedestal 78 in a push tube 79. Both distal and proximalaxial radiation baffles (not shown) may then be installed around thecrucible with the tube around the crucible and axial shields. The radialshields can then be loaded along with the pedestal mount. The crucibleis preferably positioned with respect to the radial shield geometry suchthat the nucleation tip is below the location of the sharp gradientcreated by the reflective metallic shields, and the tip initially ismaintained at a higher temperature than the source material so thatlittle or no nucleation occurs during a warm-up.

The growth chamber is then closed and evacuated, as described above, toreduce trace atmosphere contamination of the nucleation process andresulting AlN boule. In various versions of these embodiments, followingevacuation, e.g. to less than about 1 Pa employing a mechanical Welchpump with minimum pressure of about −0.5 Pa, the chamber is filled witha forming gas blend of 3% H₂ and 97% N₂ to a pressure of about 100 kPaand then evacuated again to less than 10 mTorr. This refill and pumpprocess can be carried out three times or more to reduce chambercontamination. Following the pump and refill processes, the chamber isfilled with the forming gas to a pressure of 117 kPa. High-purity gradegas available from Awesco (99.999% certified) can be used to furtherensure a clean growth chamber atmosphere.

During a ramp to the growth temperature, the pressure in the chamberincreases until the target growth pressure of 124 kPa is reached. Afterreaching the operating pressure, the chamber pressure can beperiodically checked and incrementally adjusted by releasing gas fromthe chamber to a vent line in order to keep the chamber pressure between124 kPA and 125 kPa.

In some embodiments, the power supply is an RF oscillator with a maximumpower output of 40 kW at 450 kHz. The power output of this machine ispreferably controlled with saturable reactors driven through a GPIBinterface with a Keithly ABC 0-3 amp DC current supply. The growthtemperature inside the furnace can be increased in two ramp segments.For example, the first segment of the ramp can be linear for about 15minutes taking the top axial optical pyrometer temperature to about1800° C. The second ramp segment can then be linear for approximately 4hours taking the top axial temperature to about 2200° C. The chamber canbe then maintained at growth temperature and pressure for a period ofabout 1 hour. Then, the crucible is moved up by the push rod at a rateof, for example, approximately 0.63 mm/hr. During the growth run, thispush rate is held constant, such that the total travel is about 32 mm,producing a single-crystal AlN boule that reached about 35 mm in lengthand about 11 mm in diameter. Shorter or longer crystals can be producedby varying the travel distance (which is directly related to the pushtime). Following completion of the vertical travel of the crucible, thecurrent through the saturable reactors of the oscillator is ramped down(which decreases the power to the furnace), e.g. from 1.4 A to 0.7 A in20 hrs. Once the saturable reactor current reaches 0.7 A, it can bequickly reduced to 0. Once the system is at room temperature, thechamber can be pumped to less than 1 Pa and backfilled to atmosphericpressure with the forming gas, allowing the chamber to be opened and thegrowth crucible assembly removed from the furnace for evaluation.

In the embodiments of the invention described above, a crucible tiphaving a narrow tip is employed so as to minimize the number of nucleithat form in the tip during the nucleation stage. In some versions, asingle nucleus fills the narrow tip, such that a single crystal forms inthe tip. In addition, in many embodiments, the orientation of thecrystal is close to having its c-axis along the growth direction.

It has been experimentally confirmed that the single-crystal AlNmaterial produced in accordance with implementations of the embodimentsdescribed above to be of very high quality. Total dislocation densitieswere found to average 1,000 cm⁻², with some regions of the singlecrystal having between 100 and 500 dislocation per cm² and otherregions, including those exceeding 5×5 mm², having no visibledislocations at all. Triple-crystal x-ray rocking curve widths wereobserved to be less than 15 arcsec and were typically limited byinstrument resolution.

Seeded-Crystal Growth Using an Encased Seed

In some embodiments, single-crystal AlN boules up to about 50 mm indiameter are grown by employing an encased AlN seed prepared by theself-seeded growth process described above. In these embodiments, thetip of the conical crucible used in the self-seeded process is cut offonce it is filled with an AlN single crystal. In many versions of theseembodiments, the tip is cut off to a length of approximately 15 to 20mm. At this length, the cross-section of the AlN single crystal isapproximately 5 to 7 mm in diameter.

In these embodiments, the crucible is fabricated and assembled asdescribed above in connection with FIGS. 7-8 with the cut-off tip of theself-seeded growth crucible filled with a single crystal of AlN, i.e. anencased AlN boule, disposed in a top portion of the crucible. Similar tothe embodiments described above, a source of high purity AlN, e.g.having an impurity concentration lower than 500 ppm by weight, is placedin a bottom portion of the crucible. As described above, the AlNmaterial in the source base is preferably formed from very high-puritychunks of AlN ceramic prepared by reacting pure Al with pure nitrogenand using sublimation transport to achieve near theoretical density ofAlN ceramic in the source cup. During fabrication, the single crystal ofAlN grows on the surface of the encased AlN seed and expands in diameterinto a first expansion zone A and then a second expansion zone B of thecrucible.

The furnace used to grow the single crystal self-seeded materialdescribed above may not be large enough to grow crystals larger thanabout 20 mm in diameter. Thus, for the larger-diameter crystals, a Hamcofurnace coupled to a 450 kHz, 100 kW power supply can be employed. Thearrangement of the RF coil and the radiation shield is the same asdescribed above.

In various embodiments of the invention, the crucible shown in FIG. 7 isplaced into the furnace in a manner discussed above in connection withFIG. 2. At this point, both distal and proximal axial radiation bafflesare installed around the crucible with the tube around the crucible andaxial shields. The radial shields are then loaded along with a pedestalmount. The crucible is preferably positioned with respect to the radialshield geometry such that a seed equilibrium position (SEP) wasachieved. This means that once the furnace is heated to the growthtemperature, the seed crystal will neither grow nor shrink while it ispositioned at the SEP. The SEP can be determined experimentally anddepends on the growth temperature used.

The growth chamber is then closed and evacuated as described above toreduce trace atmosphere contamination of the growth cell, nucleationprocess and resulting AlN boule. In particular versions of theseembodiments, following pump-down to less than 7 mPa, e.g. using a turbopump with a minimum pressure of about 0.4 mPa, the chamber is filledwith a forming gas blend of 3% H₂ and 97% N₂ to a pressure of about 122kPa. Following the pump and refill process, the chamber is filled withthe forming gas for the start of the growth process to a pressure of 117kPa. As described above, a high-purity grade gas available from Awesco(99.999% certified) can be used to further ensure a clean growth chamberatmosphere.

During a ramp to the growth temperature, the pressure in the chamberincreases until the target growth is reached. After reaching theoperating pressure, the chamber pressure can be periodically checked andincrementally adjusted by releasing gas from the chamber to a vent linein order to keep the chamber pressure between 124 kPa and 125 kPa.

The growth temperature inside the furnace can be increased in twosegments. For example, in the first segment, the temperature is linearlyincreased from the room temperature to about 1800° C. in 30 minutes.Then, the second ramp segment to the final growth temperature determinedby the optical pyrometer, e.g. for three hours, can be initiated afteroperator inspection. The growth temperature can be within 50° C. of thetungsten eutectic temperature (in the presence of AlN at 124 kPa totalpressure). This eutectic temperature can be used to calibrate theoptical pyrometer used to control the temperature during the growth run.

The chamber is then maintained at the growth temperature and pressurefor a period of, for example, 1 hour. The vertical drive then pushes thecrucible up at a rate ranging from about 0.4 to 0.9 mm/hr, for example,at approximately 0.455 mm/hr. In a particular version, during the growthrun, this push rate is held constant and the total travel is about 35mm, producing a single crystal AlN boule that reached about 50 mm indiameter and 35 mm in length. Shorter or longer crystals can be producedby varying the distance the crucible is pushed or equivalently byvarying the push time.

Following completion of the vertical travel, the vertical motion of thecrucible is stopped and the pressure of the furnace is increased to 157kPa by adding more high-purity forming gas. The power to the furnace isthen linearly reduced to zero, for example, in 6 hrs and the system isallowed to cool to room temperature. Following the cool down, thechamber is pumped to, e.g., less than about 1 mPa and backfilled toatmospheric pressure with forming gas. The chamber is then opened andthe growth crucible assembly removed from the furnace for evaluation.

In various embodiments, after orienting the resulting AlN boule, e.g. byemploying the x-ray Laue diffraction technique, the boule is encased inepoxy, e.g. VALTRON available from Valtech, and then ground down to a25-mm diameter cylinder having its longitudinal axis oriented along thedesired crystallographic direction. Once the oriented cylinder isproduced, it is once again examined by the x-ray Laue diffractiontechnique to determine precise orientation (within +/−0.2°) and thensliced with a wire saw, e.g. the Model DT480 saw, for example, the oneavailable from Diamond Wire Technologies, into a wafer shown in FIG. 9.Those skilled in the art of semiconductor wafer preparation will readilyrecognize that there are many alternatives for slicing the AlN crystalusing diamond-coated ID and OD saws. The surface of the wafer is thenprepared for epitaxial growth utilizing, for example, one or moretechniques described in the '838 patent.

Seeded Growth Using Polished AlN Wafers

To improve the quality of the single-crystal AlN for certainapplications, it is desirable able to determine the orientation of theseed crystal by mining it from a source bulk crystal. Accordingly, insome embodiments, a piece of AlN material having a knowncrystallographic orientation is used as a seed from which bulk materialcan then be grown. In order to further increase the diameter of the bulkcrystal, in certain versions of these embodiments, several seeds aremounted next to each other such that crystallographic directions ofthese seeds are substantially aligned. Bulk crystals grown thereover mayexhibit certain mosaicity, i.e. regions with slightly differentorientations, but this effect can be minimized by choosing appropriategrowth conditions. In a particular embodiment, a polished AlN wafersliced from a bulk crystal is employed as a seed, offering a number ofbenefits, including standardization and improved control over the growthdirection.

In order to grow high-quality AlN crystals, very high temperatures, forexample exceeding 2100° C., are generally desirable. At the same time,as discussed above, high thermal gradients are needed to providesufficient mass transport from the source material to the seed crystal.If not chosen properly, these growth conditions can result inevaporation of seed material or its total destruction and loss.

Preferably, the mounting technique employed in these embodiments tosecure AlN seeds entails:

-   -   (1) employing a seed holder and/or adhesive compound that is        sufficiently strong to secure the seed and the crystal being        grown;    -   (2) protecting the opposite side of the seed during growth to        avoid re-evaporation of the AlN, as this may result in formation        of planar and/or extended void defects; and    -   (3) avoiding contamination of the crystal and the crucible by        the material chosen to protect the opposite side of the seed or        as the adhesive;

FIG. 10 schematically depicts one exemplary embodiment of the apparatusfor seeded bulk crystal growth. An AlN wafer 100 having a thickness ofapproximately 500 μm is used as a seed. The AlN wafer is sliced from abulk AlN crystal and polished utilizing, for example, one or moretechniques described in the '838 patent. This approach enables growth ofAlN single-crystal boules with a predetermined seed orientation. Inaddition, either the nitrogen polarity or aluminum polarity can be usedfor c-face seed orientations where the seed face is oriented parallel tothe (0001) plane. A seed which oriented with its face parallel to the(11 20) (the so-called a-plane) or with its face oriented parallel tothe (10 10) (the so-called m-plane) will allow the growth of boules fromwhich it will be easier (with less material waste) to cut substrateswith these special orientations. The a-plane and m-plane substrates willhave non-polar surfaces which are particularly desired for various kindsof electronic and opto-electronic devices which can have superiorproperties over devices grown on polar substrates. An example of asuperior device property includes the elimination of the spontaneouspolarization and piezo-electric fields in the quantum wells used forlight emitting diodes (LEDs) and laser diodes (LDs). Another example, isthe enhanced impact ionization anisotropy for electrons and holes movingalong the nitride semiconductor basal plane which is desirable foravalanche photo diodes and would be facilitated by a-plane or m-planesubstrates.

In some embodiments, the AlN seeded bulk crystal growth is carried outin the tungsten crucible 110 using a high-purity AlN source 120, forexample, as described above in connection with FIGS. 7-9. The tungstencrucible is placed into an inductively heated furnace, as describedabove, so that the temperature gradient between the source and the seedmaterial drives vapor species to move from hotter high purity AlNceramic source to the cooler seed crystal. The temperature at the seedinterface and the temperature gradients are monitored and carefullyadjusted, if necessary, in order to nucleate high-qualitymono-crystalline material on the seed and not destroy the AlN seed.

Preferably, the AlN seed is mounted onto a holder plate 130 using anAlN-based adhesive 140. Using an adhesive for seed mounting generallysimplifies the fabrication procedure and improves the seed backprotection and mechanical adhesion. The AlN ceramic adhesive may containat least 75% AlN ceramic and silicate solution that provides adhesiveproperties. One suitable example of such an adhesive is Ceramabond-865available from Aremco Products, Inc.

In a particular version, the AlN seed is mounted using followingprocedure:

-   -   (1) AlN adhesive is mixed and applied to the holder plate using        a brush to a thickness not exceeding about 0.2 mm;    -   (2) The AlN seed is placed on the adhesive; and then    -   (3) the holder plate along with the seed is placed in vacuum        chamber for about 12 hrs and then heated up to 95° C. for about        2 hrs.

Alternatively, the AlN seed can be mounted on a thin foil of Al on theholder plate. The Al is melted as the temperature of the furnace israised above 630° C. (the melting point of Al), wetting the back of theseed and the holder plate. As the temperature is raised further, the Alwill react with N₂ in the furnace to form AlN and will secure the seedto the holder plate. This technique may require that the AlN seed beheld in place (either by gravity or mechanically) until a sufficientamount of the Al has reacted to form AlN which will secure the seed tothe holder without further need of mechanical support.

Referring to FIG. 11A, in one particular embodiment, a predominantlysingle-crystal boule having a diameter of about 15 mm was grown using apolished AlN wafer sliced from a bulk crystal as a seed as describedabove and then sliced into wafers approximately 1 mm thick. The firstslice (at the left) shows the original seed crystal (which wasapproximately 10 mm in diameter) in the center. This seed crystalpropagated through the boule as seen in the center sections ofsubsequent slices. The boule can be made single crystal by using a seedcrystal of the same diameter as the desired boule. Larger seed crystalscan be produced using the techniques described above. In anotherembodiment, multiple seed crystals are mounted and carefully oriented inthe same direction to increase the size of seeded region.

In yet another embodiment, the size of the seeded region has beenexpanded during growth as described below. The seed expansion has beendone starting with smaller seeds, with faces parallel to the (0001)plane, mounted on a tungsten plate in a furnace designed to grow boulesthat are 2.4″ (61 mm) in diameter. Under proper growth conditions, ithas been found that the seeded region of the resulting boule can beincreased as shown in FIG. 11B. The (0001) orientation (the so-calledc-face) was used because growth in the a-b plane perpendicular to thec-axis is expected to be more rapid than along the c-axis, making ac-face seed easier to expand. The proper conditions are a combination ofseveral additional factors including the self-seeded nucleation barrierand the use of thermal gradient factors to aid in the seed crystal areaexpansion with increasing growth distance. Brief descriptions of both ofthese conditions are given below.

Self-seeded nucleation barrier: There is a nucleation barrier to thegrowth of AlN on tungsten. That is, the vapor above the tungstencrucible will tend to be supersaturated unless AlN nuclei are availablefor growth. Some embodiments of the present invention takes advantage ofthis by having a seeded region take up some part of the full diameterseed mounting plate which is surrounded by an unseeded, bare region.Since adsorption of aluminum and nitrogen from the vapor onto the seedis favored over deposition on the bare crucible wall, the seed isfavored to expand laterally in favor of creating new self-seededcritical nuclei next to the seed. Under properly controlled conditionsthis process can be used to increase the seeded area per growth cycle.

Thermal gradient factors: The process of crystal growth requires heatextraction which is controlled by arrangements of insulators/heaters inthe system. Properly arranging insulations so that the seed is thecoolest part of the upper crucible and cooler than the source duringgrowth is important to the process. Further tailoring this insulationwhen using a small seed to be expanded during the growth aids inexpansion of the seed by making the seed cooler than the unseededlateral region. This thermal arrangement makes self-seeded nucleation'sneighboring the seed less favored by limiting heat extraction. As thecrystal grows at high temperature and with sufficient source material,given sufficient time to reach an equilibrium point during the growthrun the interface of the crystal will follow the isotherms of the system(insulation/heaters etc). The proper interface shape to favor seedexpansion is slightly convex in the growth direction—the curvature ofthe gradient will aid expansion.

The plot depicted in FIG. 11B shows a collection of data documentingthis behavior, the legend on the chart referencing the boule number orproduction number of the growth run. The significance of this expansion,along with a high quality of the grown material, is that a small seedsequentially grown out by this process generally yields larger, higherquality seeds (or substrates) in each process cycle. Reprocessing andreseeding with this expanded material is one path to achieving largediameter seeded growth.

Referring to FIG. 11C, in another particular embodiment of theinvention, a two-inch diameter AlN substrate, approximately 0.5 mmthick, was cut from a 2.4″ boule grown as described above. The 2.4″boule was ground down to 2″ diameter before slicing.

Defect Density in Pseudomorphic Epitaxial Layers

As discussed above, the AlN boules grown by the techniques disclosedherein have very low dislocation densities—under 10,000 cm⁻², typicallyabout 1,000 cm⁻², and, in certain embodiments, under 500 cm⁻² and evenunder 100 cm⁻²—enabling fabrication of highly-efficient UV LEDs and LDs,as well as electronic devices, such as transistors for high frequency(>2 GHz), high power operation. However, the lattice mismatch betweenpure AlN and pure GaN of about 2.3% may result in a high threadingdislocation density if thick layers of either GaN or low-Al-contentAl_(x)Ga_(1-x)N layers are grown on the low dislocation densitysubstrate. However, thin pseudomorphic GaN and Al_(x)Ga_(1-x)N can begrown with significantly reduced defect densities. As a result,pseudomorphic devices grown on closely-matched AlN substrates mayinclude epitaxial layers with levels of defect densities similar tothose of the underlying substrate. Thus, use of thin, pseudomorphiclayers broadens the range of low defect device structures significantly,allowing for enhanced performance and increase device reliability inmany new nitride devices, as discussed in more detail below.

In various embodiments, thin GaN and Al_(x)Ga_(1-x)N epitaxial films aregrown pseudomorphically, eliminating the formation of misfitdislocations and resulting in an overall reduction of defects that arecommonly observed in these films grown on foreign substrates such assapphire, SiC and silicon. The key is defect reduction by keeping theentire structure lattice-matched to AlN lattice parameter perpendicularto the growth direction. Since GaN and Al_(x)Ga_(1-x)N layers haveintrinsically larger lattice parameter, this can only be accomplished bystraining the epitaxial layer. In a so-called pseudomorphic layer, thestrain in the plane parallel to the substrate surface is exactly equalto the lattice mismatch between the AlN substrate and intrinsic latticeparameter of the epitaxial layer. A schematic cross-section of oneexemplary epitaxial profile is shown in FIG. 12, indicating thepseudomorphic layers fabricated in accordance various embodiments of theinvention.

From simple calculations where the strain energy in the pseudomorphiclayer is balanced against the extra energy involved in creatingdislocations, the critical layer thickness for different latticemismatches (which, in this case, is controlled by the percentage of GaNin the epitaxial layer) can be calculated. (See, for example, Matthewsand Blakeslee in the J. Crystal Growth 27, 118 (1974), and U.S. Pat. No.4,088,515). The AlGaN supply layers (−60% Al) can be 40 nm thick andremain pseudomorphic. At 50%, the critical thickness goes down to 30 nm.The critical thickness of pure GaN is estimated to be 12.5 nm withrespect to AlN. In reality, slightly thicker pseudomorphic layers cantypically be grown because the formation of strain-relieving misfitdislocations is kinetically hindered. In addition, putting one or moreadditional AlN epitaxial layers at the top of the device structure as abuffer may also help stabilize the pseudomorphic layers with lowdislocation densities.

Two different device structures can be considered that have differentphysical properties in addition to the defect density reduction aspreviously described. The first kind of structure is pseudomorphic andgrown on non-polar faces, including a- and m-plane crystal orientationsas defined by the AlN substrate. This minimizes the polarization effectsin the crystal, makes the device inherently insensitive to polarizationeffects associated with the surface, as well as eliminates or minimizesDC to RF dispersion observed on conventional devices grown along thec-plane. On the other hand, pseudomorphic structures grown on thec-plane along the [0001] direction have strong polarization effectswhich influence the charge distribution within the device. Preferably,the polarization charge at the channel/barrier interface is carefullyincreased to counteract backside depletion effects associated with theAlN/GaN hetero-interface transitioning from the AlN buffer structure.

Defect Density in Non-Pseudomorphic Epitaxial Layers

For some device structures, it is desirable to have a relatively thick,GaN epitaxial layer on the AlN substrate (where the lattice parameter ofthe GaN is approximately 2.4% larger than that of the AlN). Otherapplications may require a relatively thick, epitaxial layer consistingof a binary or tertiary alloy combination selected from the groupconsisting of AlN, GaN, and InN. These thick epitaxial structures aredesired because of the high thermal conductivity of the AlN substrateand the very small total thermal expansion mismatch between AlN, GaN andalloys of these two materials. For example, a thick GaN epitaxial layeron AlN with a low threading dislocation density is desired forfabrication of 405 nm laser diodes for the Blu-ray DVD. It is alsodesired for fabrication of efficient blue and white light emittingdiodes (LEDs).

As the thickness of a pseudomorphically strained alloy is increased, theenergy associated with the strain also grows and will need to berelieved by some mechanism once the thickness exceeds some criticalvalue as described above. A common mechanism for strain relieve inlattice mismatched epitaxial layers is the formation of mismatchdislocations. While the strain relieving part of the misfit dislocationwill run parallel or nearly parallel to the interface between theepitaxial layer and the AlN substrate, the mechanisms for introducingthe misfit dislocations into the layer will nearly always result in twothreading components which will run through the epitaxial layer to thesurface. This is because the dislocation line must either end on asurface or some other defect. For a given type of dislocation (which isdetermined by the Burgers vector of the dislocation and the orientationof the dislocation line), the density of misfit dislocations runningparallel to the substrate interface will determine the amount of latticemismatch which has been relieved by the misfit dislocations. However,while the density of threading dislocations will generally increase asthe density of misfit dislocations parallel to the interface increases,the threading density will decrease as the average length that themisfit dislocation runs parallel to the interface is increased.Unfortunately, as the density of threading dislocations increase, theaverage length of misfit dislocations is generally observed to decrease.This phenomena is generally attributed to the entanglement of thethreading dislocations which tends to impede their ability to moveparallel to the interface and increase the length of the misfitdislocation.

A commonly used way to prevent the entanglement of dislocations andincrease the length of the misfit dislocations is to gradually grade thealloy composition so as to cause the lattice mismatch to increaseslowly. Once the critical layer thickness is exceeded, it is desirableto keep the density of threading dislocations low so that they (thethreading components of the dislocations) can move over relatively largedistances and be connected by long lengths of misfit dislocations. Inthis way, lattice-mismatch strain can be nearly completely removed inthe growth of a thick alloy layer by having a gradual transition layerbetween the substrate and the desired alloy composition layer which isdesired for the fabrication of a particular active electronic oroptoelectronic device.

Once thicker layers exceed the pseudomorphic limit (which is determinedby the size of the lattice mismatch), they will require the formation ofmisfit dislocations to relieve strain in the epitaxial layer. However,by using appropriately graded structures, from pure AlN to the desiredalloy structure (or to GaN), the density of threading dislocations inthe final epitaxial layer can be reduced. These graded structures can begraded continuously, in a series of steps, or as a superlattice wherethe effective lattice parameter of the grading layer is alternativelyincreased and then decreased in order to impede the propagation ofthreading dislocations into the epitaxial layer. Similar approaches havebeen used to reduce the density of misfit dislocations in otherlattice-mismatched semiconductor heterostructures such as GaAs on Si(lattice mismatch of 4.2%) and Ge on Si (lattice mismatch of 4%) tolevels below 10⁶ cm⁻². However, it is crucial that the initial crystalsubstrate (Si in the above-referenced example) have much lowerdislocation density. Accordingly, high-quality AlN substrates describedherein with initial dislocation densities below 10,000 cm⁻² facilitategrowth of thick GaN epitaxial layers thereon where the average threadingdislocation density within the GaN is less than about 10⁶ cm⁻². Otherthick semiconductor structures (which exceed the pseudomorphic limit) ofbinary and ternary alloys of AlN, GaN and InN can also be grown on thelow dislocation density AlN substrates with similar results when theappropriate graded layer structure is used.

While the general approach of reducing threading dislocations in anepitaxial exceeding the critical layer thickness through the use ofgraded intermediary layers is understood, it is difficult to predict howsuccessful this approach will be for any particular semiconductor systembecause of the material-dependent interaction between dislocations.Generally it is best to make the grading layers thick so that the rateof change in lattice parameter is kept small. For instance, a 1% changein lattice parameter may be graded over a thickness greater than 1 μm.Because it is impractical to grow very thick layers, however, steppedstructures or superlattices are often used to reduce the thickness ofthe intermediary layer while preserving a low threading dislocationdensity. In addition, it is generally very difficult to reduce thethreading dislocation below 10⁶ cm⁻² for lattice mismatches exceeding acouple of percent. Thus, it was surprising to discover that the densityof threading dislocations in an epitaxial layer of GaN on a lowdislocation density AlN substrate could be reduced to less than 5×10⁵cm⁻² with the use of a linearly graded buffer layer that was onlyapproximately 0.1 μm thick.

For example, a 1-μm-thick GaN layer with low dislocation density (lessthan 4×10⁵ cm⁻²) may be produced on an AlN substrate through growth of agraded buffer layer. In one procedure, a large c-face AlN substrate withlow dislocation density (roughly 5×10³ cm⁻²) was prepared as describedabove. The Al-polarity surface of the c-face AlN substrate—the (0001)face—was prepared as described in U.S. Pat. No. 7,037,838 (the entiredisclosure of which is hereby incorporated by reference). Afterintroducing the substrate into a Thomas Swan OMVPE reactor, thesubstrate was heated to 1150° C. under flowing hydrogen and ammonia gasmixture. Trimethylaluminum (TMA) was then introduced and a 0.6-μm-thickAlN buffer layer was grown on the substrate at an approximate growthrate of 0.3 μm/hr. A graded layer Al_(x)Ga_(1-x)N was then grown byswitching in trimethylgallium (TMG) and ramping down the TMA gas flow tozero over a 15 minute interval to grow approximately 0.1 μm of linearlygraded alloy, then ramping up the TMG to the desired growth rate. TheTMG should be introduced at a low flow rate so as to initially producean Al_(x)Ga_(1-x)N layer with a high Al concentration. It was expectedthat the initial Al concentration should exceed 90%. However, goodresults were obtained (dislocation density<5×10⁵ cm⁻²) even when theinitial Al content was as low as 35% which was then linearly graded overa 0.1 μm thickness. Lower dislocation densities may be obtained byincreasing the grading layer thickness to over 1 μm and by increasingthe initial Al concentration in the grading layer. In addition, addingmultiple step layers or superlattices where the Al concentration isinitially stepped down and then back up (by as little as 2% and as muchas 50%) will also hinder the threading dislocation up through thegrowing epitaxial layer. These stepped layers or superlattices mayinstead cause the threading dislocation to bend so as to propagateparallel to the substrate interface and increase the length of themisfit dislocation, thereby relieving more lattice mismatch strain witha reduced threading dislocation density. These stepped layers or thethickness of a single period in a superlattice may be as thin as 1 nmand as thick as 0.1 μm. In general, increasing the number of steppedlayer or periods in the superlattice will further reduce the density ofthreading dislocations. However, practical limitations on the length ofgrowth time and the difficulty of the growth recipe makes exceeding 30stepped layers or exceeding a 30-period superlattice unattractive, asthe rate that threading dislocations are reduced generally decreases.

The dislocation density in the GaN layer grown on AlN substrates wasmeasured by using molten KOH etch. The densities of the etch pitsmeasured in Nomarski microscopy, after a five minute etch, variedbetween 9×10⁴ and 5×10⁵ cm⁻².

Although the foregoing description relates to AlN crystals and theirgrowth, the skilled artisan will recognize that various aspects of theforegoing description may be used to fabricate other crystallinesubstrates, such as ZnO, LiGaO₂, MgAl₂O₄ or Si, without departing fromthe spirit and scope of the present invention. For example, the presentmethod may be used to produce single crystals of ZnO. In thisembodiment, Zn and O₂ vapor are introduced into a crystal growthenclosure, as previously described with respect to Al and N. The O₂partial pressure in the crystal growth enclosure is held at greater thanstoichiometric pressure relative to the Zn partial pressure. Inaddition, one or more nucleation sites are contained in the crystalgrowth enclosure, and the nucleation site(s) are cooled relative toother locations in the crystal growth enclosure. The Zn and O₂ vapor isdeposited under conditions capable of growing single crystalline ZnOoriginating at the nucleation site, as described herein with respect toAlN. Optionally, the cooling step may be facilitated by placing thecrystal growth enclosure within a temperature gradient, and the crystalgrowth enclosure may then be moved through the temperature gradientwhile depositing the vapor, as discussed above.

Skilled artisans will also readily recognize that while variousembodiments of the present invention have been described herein asutilizing a seed crystal to promote crystal growth, the teachings hereinmay also be used for unseeded crystal growth, without departing from thescope and spirit of the present invention.

The modifications to the various aspects of the present inventiondescribed hereinabove are merely exemplary. It is understood that othermodifications to the illustrative embodiments will readily occur topersons with ordinary skill in the art. All such modifications andvariations are deemed to be within the scope and spirit of the presentinvention as defined by the accompanying claims. All of the patents,patent applications, journal articles, and other references mentionedabove are incorporated herein by reference.

1.-19. (canceled)
 20. A solid-state ultraviolet (UV) light-emittingdevice comprising: at least two layers each independently selected fromthe group consisting of AlN, GaN, InN, and any binary or tertiary alloycombination thereof; and therebelow, an undoped AlGaN graded layerdisposed over a substrate of AlN, a bottom portion of the AlGaN gradedlayer being lattice-matched to the substrate, wherein the substrate ofAlN has a dislocation density of about 10,000 cm⁻² or less.
 21. Thelight-emitting device of claim 20, wherein the light-emitting deviceemits UV radiation at wavelengths ranging from about 200 nm and 320 nm.22. The light-emitting device of claim 20, wherein a surface of thesubstrate has substantially a single orientation.
 23. The light-emittingdevice of claim 22, wherein the orientation is c-face, Al-polarity. 24.The light-emitting device of claim 22, wherein the orientation isc-face, N-polarity.
 25. The light-emitting device of claim 22, whereinthe orientation is m-face.
 26. The light-emitting device of claim 22,wherein the orientation is a-face.
 27. The light-emitting device ofclaim 20, wherein a grading rate of the AlGaN graded layer isapproximately 50% Ga/μm.
 28. The light-emitting device of claim 20,wherein a threading dislocation density of the at least two layers isless than a threading dislocation density of a structure comprising theat least two layers disposed over a substrate of AlN having adislocation density of about 10,000 cm⁻² or less without an AlGaN gradedlayer therebetween.
 29. The light-emitting device of claim 20, wherein athickness of the AlGaN graded layer is approximately 1 μm.
 30. Thelight-emitting device of claim 20, wherein the AlGaN graded layer isgraded from approximately pure AlN to approximately Al_(0.5)Ga_(0.5)N.31. The light-emitting device of claim 20, further comprising two metalcontacts disposed above the AlGaN graded layer and in direct contactwith the at least two layers.
 32. The light-emitting device of claim 20,wherein the light-emitting device is a light-emitting diode.
 33. Thelight-emitting device of claim 20, wherein the light-emitting device isa laser diode.
 34. The light-emitting device of claim 33, furthercomprising a plurality of quantum wells disposed between the at leasttwo layers.
 35. The light-emitting device of claim 33, wherein the atleast two layers are patterned to laterally confine radiation.
 36. Asolid-state ultraviolet (UV) light-emitting diode comprising at leasttwo layers each independently selected from the group consisting of AlN,GaN, InN, and any binary or tertiary alloy combination thereof, whereinthe light-emitting diode is disposed atop a substrate of AlN having adislocation density of about 10,000 cm⁻² or less and is configured toconvert a >10% fraction of electric power into UV radiation power. 37.The light-emitting diode of claim 36, further comprising an AlGaN gradedlayer disposed between the at least two layers and the substrate. 38.The light-emitting diode of claim 37, wherein the AlGaN graded layer isundoped.
 39. The light-emitting diode of claim 37, further comprisingtwo metal contacts disposed above the AlGaN graded layer and in directcontact with the at least two layers.
 40. The light-emitting diode ofclaim 37, wherein a grading rate of the AlGaN graded layer isapproximately 50% Ga/μm.
 41. The light-emitting diode of claim 37,wherein a threading dislocation density of the at least two layers isless than a threading dislocation density of a structure comprising theat least two layers disposed over a substrate of AlN having adislocation density of about 10,000 cm⁻² or less without an AlGaN gradedlayer therebetween.
 42. The light-emitting diode of claim 37, wherein athickness of the AlGaN graded layer is approximately 1 μm.
 43. Thelight-emitting diode of claim 37, wherein the AlGaN graded layer isgraded from approximately pure AlN to approximately Al_(0.5)Ga_(0.5)N.